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Landolt-Börnstein Numerical Data and Functional Relationships in Science and Technology New Series / Editor in Chief: W. Martienssen

Group IV: Physical Chemistry Volume 11

Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data critically evaluated by MSIT® Subvolume E Refractory Metal Systems Part 1 Selected Systems from Al-B-C to B-Hf-W Editors G. Effenberg and S. Ilyenko

Authors Materials Science and International Team, MSIT®

ISSN

1615-2018 (Physical Chemistry)

ISBN

978-3-540-88052-3 Springer Berlin Heidelberg New York

Library of Congress Cataloging in Publication Data Zahlenwerte und Funktionen aus Naturwissenschaften und Technik, Neue Serie Editor in Chief: W. Martienssen Vol. IV/11E1: Editors: G. Effenberg, S. Ilyenko At head of title: Landolt-Börnstein. Added t.p.: Numerical data and functional relationships in science and technology. Tables chiefly in English. Intended to supersede the Physikalisch-chemische Tabellen by H. Landolt and R. Börnstein of which the 6th ed. began publication in 1950 under title: Zahlenwerte und Funktionen aus Physik, Chemie, Astronomie, Geophysik und Technik. Vols. published after v. 1 of group I have imprint: Berlin, New York, Springer-Verlag Includes bibliographies. 1. Physics--Tables. 2. Chemistry--Tables. 3. Engineering--Tables. I. Börnstein, R. (Richard), 1852-1913. II. Landolt, H. (Hans), 1831-1910. III. Physikalisch-chemische Tabellen. IV. Title: Numerical data and functional relationships in science and technology. QC61.23 502'.12 62-53136 This work is subject to copyright. All rights are reserved, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilm or in other ways, and storage in data banks. Duplication of this publication or parts thereof is permitted only under the provisions of the German Copyright Law of September 9, 1965, in its current version, and permission for use must always be obtained from Springer-Verlag. Violations are liable for prosecution act under German Copyright Law. Springer is a part of Springer Science+Business Media springeronline.com © Springer-Verlag Berlin Heidelberg 2009 Printed in Germany The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Product Liability: The data and other information in this handbook have been carefully extracted and evaluated by experts from the original literature. Furthermore, they have been checked for correctness by authors and the editorial staff before printing. Nevertheless, the publisher can give no guarantee for the correctness of the data and information provided. In any individual case of application, the respective user must check the correctness by consulting other relevant sources of information. Cover layout: Erich Kirchner, Heidelberg Typesetting: Materials Science International Services GmbH, Stuttgart Printing and Binding: AZ Druck, Kempten/Allgäu

SPIN: 1091 6131

63/3020 - 5 4 3 2 1 0 – Printed on acid-free paper

Editors: Associate Editor:

Günter Effenberg Svitlana Ilyenko Oleksandr Dovbenko

MSI, Materials Science International Services GmbH Postfach 800749, D-70507, Stuttgart, Germany http://www.matport.com

Authors: Materials Science International Team, MSIT® The present series of books results from collaborative evaluation programs performed by MSI and authored by MSIT®. In this program data and knowledge are contributed by many individuals and accumulated over almost twenty years, now. The content of this volume is a subset of the ongoing MSIT® Evaluation Programs. Authors of this volume are: Zoya Alekseeva, Moscow, Russia

Jozefien De Keyzer, Heverlee, Belgium

Christian Bätzner, Stuttgart, Germany

Natalia Kol’chugina, Moscow, Russia

Natalia Bochvar, Moscow, Russia

Kostyantyn Korniyenko, Kyiv, Ukraine

Anatoliy Bondar, Kyiv, Ukraine

Artem Kozlov, Clausthal-Zellerfeld, Germany

Jörg Beuers, Hanau, Germany Gabriele Cacciamani, Genova, Italy Andriy Grytsiv, Vienna, Austria Lesley Cornish, Randburg, South Africa

Vasyl Kublii, Kyiv, Ukraine Viktor Kuznetsov, Moscow, Russia Nathalie Lebrun, Lille, France Hans Leo Lukas, Stuttgart, Germany Annelies Malfliet, Heverlee, Belgium

Damian M. Cupid, Freiberg, Germany Dmytro Pavlyuchkov, Jülich, Germany Tatiana Dobatkina, Moscow, Russia Oleksandr Dovbenko, Stuttgart, Germany Guenter Effenberg, Stuttgart, Germany Ol’ga Fabrichnaya, Freiberg, Germany Sergio Gama, Campinas, Brazil

Pierre Perrot, Lille, France Qingsheng Ran, Stuttgart, Germany Peter Rogl, Vienna, Austria Lazar Rokhlin, Moscow, Russia Eberhard E. Schmid, Frankfurt, Germany

Joachim Gröbner, Clausthal-Zellerfeld, Germany An Serbruyns, Heverlee, Belgium Mireille Harmelin, Paris, France

Elena Semenova, Kyiv, Ukraine

Frederick H. Hayes, Manchester, UK

Jean-Claude Tedenac, Montpellier, France

Michael Hoch, Cincinnati, USA

Vasyl Tomashik, Kyiv, Ukraine

Svitlana Ilyenko, Stuttgart, Germany

Mikhail Turchanin, Kramatorsk, Ukraine

Volodymyr Ivanchenko, Kyiv, Ukraine

Tamara Velikanova, Kyiv, Ukraine

Institutions The content of this volume is produced by MSI, Materials Science International Services GmbH and the international team of materials scientists, MSIT®. Contributions to this volume have been made from the following institutions:

The Baikov Institute of Metallurgy, Academy of Sciences, Moscow, Russia Degussa AG, Hanau, Germany Donbass State Mechanical Engineering Academy, Kramatorsk, Ukraine Forschungszentrum Jülich, Institut für Festkörperforschung (IFF), Institut Mikrostrukturforschung, Jülich, Germany I.M. Frantsevich Institute for Problems of Materials Science, National Academy of Sciences, Kyiv, Ukraine Institute for Semiconductor Physics, National Academy of Sciences, Kyiv, Ukraine Katholieke Universiteit Leuven, Department Metaalkunde en Toegepaste Materiaalkunde, Heverlee, Belgium G.V. Kurdyumov Institute for Metal Physics, National Academy of Sciences, Kyiv, Ukraine Materials Science International Services GmbH, Stuttgart, Germany Max-Planck-Institut für Metallforschung, Institut für Werkstoffwissenschaft, Pulvermetallurgisches Laboratorium, Stuttgart, Germany Moscow State University, Department of General Chemistry, Moscow, Russia

School of Chemical and Metallurgical Engineering, The University of the Witwatersrand, DST/NRF Centre of Excellence for Strong Material, South Afrika Technische Universität Bergakademie Freiberg, Institut für Werkstoffwissenschaft, Freiberg, Germany Technische Universität Clausthal, Metallurgisches Zentrum, Clausthal-Zellerfeld, Germany Universidade Estadual de Campinas, Instituto de Fisica “Gleb Wataghin”, DFESCM, Campinas, Brazil Universita di Genova, Dipartimento di Chimica, Genova, Italy Universität Wien, Institut für Physikalische Chemie, Wien, Austria Universite de Lille I, Laboratoire de Métallurgie Physique, Villeneuve d’ASCQ, France Universite de Montpellier II, Laboratorie de Physico-chimie de la Materiere Montpellier, France University of Cincinnati, Department of Materials Science and Engineering, Cincinnati, USA

Preface The sub-series Ternary Alloy Systems of the Landolt-Börnstein New Series provides reliable and comprehensive descriptions of the materials constitution, based on critical intellectual evaluations of all data available at the time and it critically weights the different findings, also with respect to their compatibility with today’s edge binary phase diagrams. Selected are ternary systems of importance to alloy development and systems which gained in the recent years otherwise scientific interest. In one ternary materials system, however, one may find alloys for various applications, depending on the chosen composition. Reliable phase diagrams provide scientists and engineers with basic information of eminent importance for fundamental research and for the development and optimization of materials. So collections of such diagrams are extremely useful, if the data on which they are based have been subjected to critical evaluation, like in these volumes. Critical evaluation means: there where contradictory information is published data and conclusions are being analyzed, broken down to the firm facts and re-interpreted in the light of all present knowledge. Depending on the information available this can be a very difficult task to achieve. Critical evaluations establish descriptions of reliably known phase configurations and related data. The evaluations are performed by MSIT®, Materials Science International Team, a group of scientists working together since 1984. Within this team skilled expertise is available for a broad range of methods, materials and applications. This joint competence is employed in the critical evaluation of the often conflicting literature data. Particularly helpful in this are targeted thermodynamic and atomistic calculations for individual equilibria, driving forces or complete phase diagram sections. Conclusions on phase equilibria may be drawn from direct observations e.g. by microscope, from monitoring caloric or thermal effects or measuring properties such as electric resistivity, electro-magnetic or mechanical properties. Other examples of useful methods in materials chemistry are massspectrometry, thermo-gravimetry, measurement of electro-motive forces, X-ray and microprobe analyses. In each published case the applicability of the chosen method has to be validated, the way of actually performing the experiment or computer modeling has to be validated as well and the interpretation of the results with regard to the material’s chemistry has to be verified. Therefore insight in materials constitution and phase reactions is gained from many distinctly different types of experiments, calculation and observations. Intellectual evaluations which interpret all data simultaneously reveal the chemistry of the materials system best. An additional degree of complexity is introduced by the material itself, as the state of the material under test depends heavily on its history, in particular on the way of homogenization, thermal and mechanical treatments. All this is taken into account in an MSIT® expert evaluation. To include binary data in the ternary evaluation is mandatory. Each of the three-dimensional ternary phase diagrams has edge binary systems as boundary planes; their data have to match the ternary data smoothly. At the same time each of the edge binary systems A-B is a boundary plane for many other ternary A-B-X systems. Therefore combining systematically binary and ternary evaluations increases confidence and reliability in both ternary and binary phase diagrams. This has started systematically for the first time here, by the MSIT® Evaluation Programs applied to the Landolt-Börnstein New Series. The degree of success, however, depends on both the nature of materials and scientists! The multitude of correlated or inter-dependant data requires special care. Within MSIT® an evaluation routine has been established that proceeds knowledge driven and applies both, human based expertise and electronically formatted data and software tools. MSIT® internal discussions take place in almost all evaluation works and on many different specific questions the competence of a team is added to the work of individual authors. In some cases the authors of earlier published work contributed to the knowledge

base by making their original data records available for re-interpretation. All evaluation reports published here have undergone a thorough review process in which the reviewers had access to all the original data. In publishing we have adopted a standard format that presents the reader with the data for each ternary system in a concise and consistent manner, as applied in the “MSIT® Workplace Phase Diagrams Online”. The standard format and special features of the Landolt-Börnstein compendium are explained in the Introduction to the volume. In spite of the skill and labor that have been put into this volume, it will not be faultless. All criticisms and suggestions that can help us to improve our work are very welcome. Please contact us via [email protected]. We hope that this volume will prove to be as useful for the materials scientist and engineer as the other volumes of Landolt-Börnstein New Series and the previous works of MSIT® have been. We hope that the Landolt Börnstein Sub-series, Ternary Alloy Systems will be well received by our colleagues in research and industry. On behalf of the participating authors we want to thank all those who contributed their comments and insight during the evaluation process. In particular we thank the reviewers - Pierre Perrot, Tamara Velikanova, Hans Leo Lukas, Marina Bulanova, Mikhail Turchanin, Nataliya Bochvar, Olga Fabrichnaya and Viktor Kuznetsov. We all gratefully acknowledge the dedicated scientific desk editing by Oleksandra Berezhnytska, Mariya Saltykova and Oleksandr Rogovtsov.

Günter Effenberg, Svitlana Ilyenko and Oleksandr Dovbenko

Stuttgart, June 2008

Contents IV/11E1 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume E Refractory Metal Systems Part 1 Selected Systems from Al-B-C to B-Hf-W Introduction Data Covered..........................................................................................................................................XI General ...................................................................................................................................................XI Structure of a System Report .................................................................................................................XI Introduction....................................................................................................................................XI Binary Systems ..............................................................................................................................XI Solid Phases ................................................................................................................................. XII Quasibinary Systems...................................................................................................................XIII Invariant Equilibria .....................................................................................................................XIII Liquidus, Solidus, Solvus Surfaces ............................................................................................XIII Isothermal Sections.....................................................................................................................XIII Temperature – Composition Sections ........................................................................................XIII Thermodynamics.........................................................................................................................XIII Notes on Materials Properties and Applications........................................................................XIII Miscellaneous .............................................................................................................................XIII References ...................................................................................................................................XVI General References ........................................................................................................................... XVII

Ternary Systems Al – B – C (Aluminium – Boron – Carbon) ...........................................................................................1 Al – B – Mo (Aluminium – Boron – Molybdenum) ............................................................................24 Al – B – Si (Aluminium – Boron – Silicon) .........................................................................................33 Al – C – Ti (Aluminium – Carbon – Titanium)....................................................................................41 Al – Cr – Ti (Aluminium – Chromium – Titanium).............................................................................72 Al – Fe – Nb (Aluminium – Iron – Niobium).....................................................................................109 Al – Mo – Ni (Aluminium – Molybdenum – Nickel) ........................................................................123 Al – Mo – U (Aluminium – Molybdenum – Uranium) ......................................................................144 Al – Nb – Ni (Aluminium – Niobium – Nickel).................................................................................164 Al – Nb – Si (Aluminium – Niobium – Silicon).................................................................................193 Al – Ni – V (Aluminium – Nickel – Vanadium) ................................................................................209 Al – O – Zr (Aluminium – Oxygen – Zirconium) ..............................................................................225 Al – Ta – Ti (Aluminium – Tantalum – Titanium).............................................................................242 B – C – Cr (Boron – Carbon – Chromium) ........................................................................................261 B – C – Hf (Boron – Carbon – Hafnium) ...........................................................................................282 B – C – Mo (Boron – Carbon – Molybdenum)...................................................................................306

B – C – N (Boron – Carbon – Nitrogen).............................................................................................323 B – C – Nb (Boron – Carbon – Niobium)...........................................................................................347 B – C – Si (Boron – Carbon – Silicon) ...............................................................................................367 B – C – Ta (Boron – Carbon – Tantalum) ..........................................................................................395 B – C – V (Boron – Carbon – Vanadium) ..........................................................................................408 B – C – W (Boron – Carbon – Tungsten) ...........................................................................................427 B – C – Zr (Boron – Carbon – Zirconium) .........................................................................................450 B – Cr – Mn (Boron – Chromium – Manganese) ...............................................................................476 B – Cr – Ti (Boron – Chromium – Titanium).....................................................................................485 B – Cr – Zr (Boron – Chromium – Zirconium) ..................................................................................494 B – Fe – Mo (Boron – Iron – Molybdenum) ......................................................................................500 B – Hf – Mo (Boron – Hafnium – Molybdenum) ..............................................................................514 B – Hf – W (Boron – Hafnium – Tungsten) .......................................................................................523

Introduction

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Introduction Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data

Data Covered The series focuses on light metal ternary systems and includes phase equilibria of importance for alloy development, processing or application, reporting on selected ternary systems of importance to industrial light alloy development and systems which gained otherwise scientific interest in the recent years.

General The series provides consistent phase diagram descriptions for individual ternary systems. The representation of the equilibria of ternary systems as a function of temperature results in spacial diagrams whose sections and projections are generally published in the literature. Phase equilibria are described in terms of liquidus, solidus and solvus projections, isothermal and pseudobinary sections; data on invariant equilibria are generally given in the form of tables. The world literature is thoroughly and systematically searched back to the year 1900. Then, the published data are critically evaluated by experts in materials science and reviewed. Conflicting information is commented upon and errors and inconsistencies removed wherever possible. It considers those, and only those data, which are firmly established, comments on questionable findings and justifies re-interpretations made by the authors of the evaluation reports. In general, the approach used to discuss the phase relationships is to consider changes in state and phase reactions which occur with decreasing temperature. This has influenced the terminology employed and is reflected in the tables and the reaction schemes presented. The system reports present concise descriptions and hence do not repeat in the text facts which can clearly be read from the diagrams. For most purposes the use of the compendium is expected to be self-sufficient. However, a detailed bibliography of all cited references is given to enable original sources of information to be studied if required.

Structure of a System Report The constitutional description of an alloy system consists of text and a table/diagram section which are separated by the bibliography referring to the original literature (see Fig. 1). The tables and diagrams carry the essential constitutional information and are commented on in the text if necessary. Where published data allow, the following sections are provided in each report: Landolt‐Bo¨rnstein New Series IV/11E1

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. Fig. 1 Structure of a system report

Introduction The opening text reviews briefly the status of knowledge published on the system and outlines the experimental methods that have been applied. Furthermore, attention may be drawn to questions which are still open or to cases where conclusions from the evaluation work modified the published phase diagram.

Binary Systems Where binary systems are accepted from standard compilations reference is made to these compilations. In other cases the accepted binary phase diagrams are reproduced for the convenience of the reader. The selection of the binary systems used as a basis for the evaluation of the ternary system was at the discretion of the assessor. DOI: 10.1007/978-3-540-88053-0_1 ß Springer 2009

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Solid Phases The tabular listing of solid phases incorporates knowledge of the phases which is necessary or helpful for understanding the text and diagrams. Throughout a system report a unique phase name and abbreviation is allocated to each phase. Phases with the same formulae but different space lattices (e.g. allotropic transformation) are distinguished by:

– small letters (h), high temperature modification (h2 > h1) (r), room temperature modification (1), low temperature modification (l1 > l2) – Greek letters, e.g., ε, ε´ – Roman numerals, e.g., (I) and (II) for different pressure modifications. In the table “Solid Phases” ternary phases are denoted by * and different phases are separated by horizontal lines.

Quasibinary Systems Quasibinary (pseudobinary) sections describe equilibria and can be read in the same way as binary diagrams. The notation used in quasibinary systems is the same as that of vertical sections, which are reported under “Temperature – Composition Sections”.

Invariant Equilibria The invariant equilibria of a system are listed in the table “Invariant Equilibria” and, where possible, are described by a constitutional “Reaction Scheme” (Fig. 2). The sequential numbering of invariant equilibria increases with decreasing temperature, one numbering for all binaries together and one for the ternary system. Equilibria notations are used to indicate the reactions by which phases will be

– decomposed (e- and E-type reactions) – formed (p- and P-type reactions) – transformed (U-type reactions) ¨ bergangsreaktion) is used in order to reserve the For transition reactions the letter U (U letter T to denote temperature. The letters d and D indicate degenerate equilibria which do not allow a distinction according to the above classes.

Liquidus, Solidus, Solvus Surfaces The phase equilibria are commonly shown in triangular coordinates which allow a reading of the concentration of the constituents in at.%. In some cases mass% scaling is used for better data readability (see Figs. 3 and 4). Landolt‐Bo¨rnstein New Series IV/11E1

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. Fig. 2 Typical reaction scheme

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In the polythermal projection of the liquidus surface, monovariant liquidus grooves separate phase regions of primary crystallization and, where available, isothermal lines contour the liquidus surface (see Fig. 3).

Isothermal Sections Phase equilibria at constant temperatures are plotted in the form of isothermal sections (see Fig. 4).

Temperature – Composition Sections Non-quasibinary T-x sections (or vertical sections, isopleths, polythermal sections) show the phase fields where generally the tie lines are not in the same plane as the section. The notation employed for the latter (see Fig. 5) is the same as that used for binary and pseudobinary phase diagrams.

. Fig. 3 Hypothetical liquidus surface showing notation employed


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. Fig. 4 Hypothetical isothermal section showing notation employed

Thermodynamics Experimental ternary data are reported in some system reports and reference to thermodynamic modelling is made.

Notes on Materials Properties and Applications Noteworthy physical and chemical materials properties and application areas are briefly reported if they were given in the original constitutional and phase diagram literature.

Miscellaneous In this section noteworthy features are reported which are not described in preceding paragraphs. These include graphical data not covered by the general report format, such as lattice spacing – composition data, p-T-x diagrams, etc. DOI: 10.1007/978-3-540-88053-0_1 ß Springer 2009

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. Fig. 5 Hypothetical vertical section showing notation employed

References The publications which form the bases of the assessments are listed in the following manner: [1974Hay] Hayashi, M., Azakami, T., Kamed, M., “Effects of Third Elements on the Activity of Lead in Liquid Copper Base Alloys” (in Japanese), Nippon Kogyo Kaishi, 90, 51–56 (1974) (Experimental, Thermodyn., 16) This paper, for example, whose title is given in English, is actually written in Japanese. It was published in 1974 on pages 51- 56, volume 90 of Nippon Kogyo Kaishi, the Journal of the Mining and Metallurgical Institute of Japan. It reports on experimental work that leads to thermodynamic data and it refers to 16 cross-references. Additional conventions used in citing are: # to indicate the source of accepted phase diagrams * to indicate key papers that significantly contributed to the understanding of the system. Standard reference works given in the list “General References” are cited using their abbreviations and are not included in the reference list of each individual system.


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General References [C.A.] [Curr.Cont.] [E] [G] [H] [L-B]

[Mas] [Mas2] [P] [S] [V-C] [V-C2]

Chemical Abstracts - pathways to published research in the world’s journal and patent literature http://www.cas.org/ Current Contents - bibliographic multidisciplinary current awareness Web resource - http://www.isinet. com/products/cap/ccc/ Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York (1965) Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York (1958) Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P., Kandler, H. and Stegherr, A., Structure Data of Elements and Intermetallic Phases (1971); Vol. 7, Pies, W. and Weiss, A., Crystal Structure of Inorganic Compounds, Part c, Key Elements: N, P, As, Sb, Bi, C (1979); Group 4: Macroscopic and Technical Properties of Matter, Vol. 5, Predel, B., Phase Equilibria, Crystallographic and Thermodynamic Data of Binary Alloys, Subvol. a: Ac-Au … Au-Zr (1991); Springer-Verlag, Berlin. Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys, Pergamon Press, New York, Vol. 1 (1958), Vol. 2 (1967) Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York (1969) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

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Index of Alloy Systems

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Index of Alloy Systems Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data

Index of Ternary Refractory Metal Alloy Systems Al-B-C to B-Hf-W

Al-B-C (Aluminium - Boron - Carbon) Al-B-Mo (Aluminium - Boron - Molybdenum) Al-B-Si (Aluminium - Boron - Silicon) Al-C-Ti (Aluminium - Carbon - Titanium) Al-Cr-Ti (Aluminium - Chromium - Titanium) Al-Fe-Nb (Aluminium - Iron - Niobium) Al-Mo-Ni (Aluminium - Molybdenum - Nickel) Al-Mo-U (Aluminium - Molybdenum - Uranium) Al-Nb-Ni (Aluminium - Niobium - Nickel) Al-Nb-Si (Aluminium - Niobium - Silicon) Al-Ni-V (Aluminium - Nickel - Vanadium) Al-O-Zr (Aluminium - Oxygen - Zirconium) Al-Ta-Ti (Aluminium - Tantalum - Titanium) B-C-Cr (Boron - Carbon - Chromium) B-C-Hf (Boron - Carbon - Hafnium) B-C-Mo (Boron - Carbon - Molybdenum) B-C-N (Boron - Carbon - Nitrogen) B-C-Nb (Boron - Carbon - Niobium) B-C-Si (Boron - Carbon - Silicon) B-C-Ta (Boron - Carbon - Tantalum) B-C-V (Boron - Carbon - Vanadium) B-C-W (Boron - Carbon - Tungsten) B-C-Zr (Boron - Carbon - Zirconium) B-Cr-Mn (Boron - Chromium - Manganese) B-Cr-Ti (Boron - Chromium - Titanium) B-Cr-Zr (Boron - Chromium - Zirconium) B-Fe-Mo (Boron - Iron - Molybdenum) B-Hf-Mo (Boron - Hafnium - Molybdenum) B-Hf-W (Boron - Hafnium - Tungsten)


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Aluminium – Boron – Carbon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Andriy Grytsiv, Peter Rogl

Introduction Due to the presence of carbon contaminant in aluminium borides the data on the constitution of the Al-B-C ternary system and those of the binary Al-B system have to be reviewed carefully, which was done by [1994Dus] to strictly differentiate between true aluminium borides and aluminium boron carbides. High hardness in combination with high neutron absorption, high wear resistance and impact resistance have triggered an early interest in high-strength and low-weight Al-B4C composite materials or cermets, either in bulk form with a metal binder or by reinforcing an Al-base matrix with boron carbide particles or with boron carbide-coated fibres. Understanding the phase equilibria proved of major importance in the processing of Al-B4C composites particular in finding processing criteria at temperatures high enough to promote wetting and low enough to control reactions and design microstructures. Despite much effort was spent on the synthesis and crystallographic characterization of the various ternary aluminum boron carbide compounds [1964Mat, 1965Eco, 1965Mat, 1966Gie, 1966Lip, 1969Per, 1969Wil, 1970Nei, 1977Mat, 1987Sar, 1980Ino, 1990Oka, 1992Via, 1994Kud, 1995Osc, 1996Hil1, 1996Hil2, 1997Mey], information on the equilibrium phase relations in the Al-B-C ternary system is scarce [1993Bau, 1997Via, 1998Rog]. These informations comprise early calculations of the phase equilibria disregarding the boron-rich compounds or rather assuming the phases, AlB40C4 and Al2.1B51C8, to be part of the solid solution range of “B4C” [1982Doe, 1993Kau]. Some confusion in the early experimental work on aluminum borides arose from the fact that due to contamination either from high carbon level boron starting material or from the use of graphite crucibles and substrates aluminum boron carbides were produced rather than binary aluminum borides. This is particularly true for “AlB10” [1963Wil] - shown to be “AlB24C4” or more precisely Al2.1B51C8 [1964Mat, 1967Wil, 1969Wil, 1969Per, 1990Oka] - and βAlB12 [1960Koh], later shown to be Al3B48C2 [1965Mat]. According to structural and DTA investigations [1996Hil1], Al3B48C2 exists in a tetragonal high temperature modification, which on cooling below 650˚C transforms into a bodycentered orthorhombic low temperature phase with a unique structure type. The mixture of two orthorhombic phases with coherent boundary and commensurable lattice parameters (modifications A and B), as claimed by [1965Mat, 1986Pes], thus simply explains by multiple twinning on cooling [1996Hil1]. An experimental study of the isothermal section at 1400˚C by [1993Bau] confirmed the existence of four ternary compounds Al2.1B51C8 [1964Mat, 1967Wil, 1969Wil, 1969Per], AlB40C4 [1966Gie, 1966Lip, 1970Nei], Al3B48C2 [1996Hil1] and Al3BC3 [1996Hil2]. The latter compound was first mentioned as “Al4B1–3C4” [1964Mat] and later labelled as “Al8B4C7” [1980Ino] from a cursory investigation of its crystal symmetry with X-ray single crystal Landolt‐Bo¨rnstein New Series IV/11E1

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photographs, although no details of the crystal structure were derived. The relation to a wurtzite structure was discussed [1995Osc]. A structure determination is due to [1996Hil2]. A fifth compound, Al3BC [1992Via, 1993Gon, 1997Mey, 2002Zhe], (earlier “Al4BC” [1987Sar, 1989Hal, 1990Pyz]) was reported to exist below 1000˚C [1992Via], however, was shown in the isothermal section at 1000˚C [1997Via]. From a detailed analysis (XPD, LOM, SEM, EPMA) of the Al-rich corner [1997Via] on about 30 specimens prepared from cold pressed and sintered powder compacts in the temperature region from 627 to 1000˚C, an isothermal section at 1000˚C and a tentative liquidus projection was derived assisted by a series of isothermal diffusion experiments by heating together in an alumina boat an Al-B rod and an Al-C rod. Phase equilibria at 900˚C in the Al-C rich part of the ternary Al-B-C system were established [2002Zhe] from XPD of about 45 ternary and binary alloys. Equilibrium conditions were not reached for boron-rich samples. An attempt to obtain equilibrated samples from mixtures B4C+AlB2, B4C+Al and B4C+Al4C3 were also unsuccessful. Al3BC and Al4C3 phases form very easily and are observed in all samples even after short time sintering in contrast to Al3BC3, which forms very slowly at 900˚C. On the other hand Al3BC3 was always observed in arc melted samples containing 40–60 at.% Al and 10–30 at.% B. Experimental techniques for preparation concerned (a) melting of B4C in excess of Al for the synthesis of AlB40C4 (at 1550˚C, [1970Nei]), (b) melting of boron with excess of Al in a graphite crucible for synthesis of Al3B48C2 (at 1400˚C, [1964Mat]) (c) vapor deposition at 1400 to 1600˚C for single crystals of Al3B48C2 [1967Bli] (d) hot pressing of B4C+Al in graphite dies for synthesis of Al2.1B51C8 (at 1800˚C [1966Gie, 1966Lip]), (e) infiltration of B4C by liquid Al at 1100˚C and anneal at 1000˚C to obtain Al3BC [1987Sar], (f) reaction sintering of Al+B+C powder compacts on alumina boats in sealed silica capsules 627 to 1000˚C for the synthesis of Al3BC and phase relations at 1000˚C [1992Via, 1997Via] or at 1400˚C for 10 h for the production of the single crystals of Al3B48C2 [1994Kud] (g) melting of an Al8BC mixture in alumina under argon for 160 h at 850˚C and subsequent cooling at 150K/h to RT for production of black-bluish single crystals of Al3BC [1997Mey] (h) melting of an Al40B2C3 mixture in alumina under argon at 1500˚C and subsequent cooling at 10K/h to 600˚C for production of single crystals of Al3BC3 in the form of yellow, transparent platelets [1996Hil2] and (i) Al-flux solvent method for a general production of single crystals (see i.e. [1986Kis, 1990Oka, 1996Hil1]). Samples used for the isothermal section at 1400˚C were prepared from cold compacted powder mixtures of AlB2, B4C, B and/or C, which were reaction-sintered under Ar in closed Knudsen-type graphite reactors at 1600˚C for 1h prior to 48 h heat treatment at 1400˚C [1993Bau]. Phase relations at 900˚C were studied [2002Zhe] on elemental powder compacts sintered in alumina crucibles (binary Al-B alloys) or in closed graphite crucibles (ternary alloys). The specimens were sealed in evacuated quartz ampoules and were slowly heated for 10˚/h to 720˚C (slightly above the melting point of aluminium) and kept at this temperature for 48h. After temperature was increased to 900˚C at a rate of 20˚/h, the tablets were sintered at this temperature for 1 week. Repeated repowderisation (under protective cyclohexane) and sintering at 900˚C were necessary to reach equilibrium conditions. Several studies dealt with the kinetics of wetting of B4C surfaces by liquid aluminum; detailed discussions can be found in the articles by [1979Pan] and [1989Hal]. Hot-pressing of B4.3C+Al powders at 1820˚C, 45 MPa under Ar (5 to 20 mass% Al) revealed the formation of the ternary B4C- related Al-boron carbides (solution of Al in B4C, and τ2) although the products were all thought to belong to the B4C-based solid solution [2000Liu]. With increasing Al-content (>5 mass% Al) the Al3BC3 phase evolved [2000Liu]. DOI: 10.1007/978-3-540-88053-0_3 ß Springer 2009

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Thermodynamic calculations of the Al-B-C system have been attempted by [1982Doe, 1993Wen, 1993Kau], however, are not fully consistent with experimental observations. Reviews on the constitution and on the crystal structures of the Al-B-C system have been presented by [1977Mat, 1990Luk, 1998Rog].

Binary Systems The binary systems B-C and Al-C are consistent with the critical assessments of [1998Kas, 1996Kas] and [2003Per], respectively. In spite of numerous data available from literature on the constitution of the Al-B phase diagram, contradictory results exist for the formation of aluminium diboride (Table 1). Figure 1a shows the various versions for the Al-rich part of the Al-B phase diagram. It should be noted, that recent experiments [2002Zhe] confirmed the formation of AlB2 at 900˚C, in contrast to data of [1997Via] suggesting peritectic formation at 892 ± 5˚C. In the present assessment we accept the temperature of 956 ± 5˚C for the invariant reaction L+AlB12ÐAlB2 as determined by [2000Hal]. The adopted Al-B phase diagram (Figs. 1b, 1c) is based on the assessment of [1994Dus]. The composition of the peritectic liquid at 0.55 at.% B has been confirmed by a recent thermodynamic assessment of [2001Fje]. AlB2 is still taken as a stoichiometric compound in spite of the suggestions of [1964Mat, 1999Bur, 2002Bur] for Al-deficiency in terms of Al0.9B2. Although the assessment of [1994Dus] concluded a peritectic formation of AlB12, L+(B)Ð AlB12 at 2050˚C, the thermodynamic calculation of [1993Wen] is based on congruently melting AlB12 (TM = 2150˚C).

Solid Phases The crystallographic information on all the binary and ternary phases pertinent to the Al-B-C system is listed in Table 2. Some controversy exists in the crystallographic characterization of the modifications reported for Al3B48C2. A single crystal study [1995Hil, 1996Hil1, 2000Mey] on an untwinned specimen revealed a tetragonal high temperature form (closely related to the structure of I-tetragonal boron), which on cooling undergoes a symmetry reduction resulting in microscopically twinned products that hitherto were indexed on the basis of two orthorhombic modifications, labeled A and B by [1965Mat]. The transformation was earlier proposed to be at ca. 850˚C [1960Koh, 1965Mat], whereas new results from DTA recorded 650˚C [1996Hil1]. The transition seems to be rather fast, as the low temperature modification is present in samples furnace-cooled from 1400˚C to room temperature [1993Bau]. A second point of controversy concerns the phases AlB40C4 and Al2.1B51C8 for which detailed crystallographic descriptions are available, however, AlB40C4 actually being isotypic with binary B4C, hitherto is not thoroughly established as a ternary phase independent from binary B4C. As the two structurally closely related phases AlB40C4 and Al2.1B51C8 generally are found together, a high and low temperature transition between them may be inferred [1993Bau]. Without further details the maximum solid solubility of Al in boroncarbide (“B13C2”, at 1950˚C) was reported to be 1 mass% Al (equivalent to 2 at.% Al in B4C) [1978Ekb]. Experiments to establish a possible homogeneous range for Al3BC3 (earlier: “Al8B4C7” [1980Ino], or “Al4B1–3C4” [1964Mat]), were carried out at 1830˚C by [1980Ino] resulting in a Landolt‐Bo¨rnstein New Series IV/11E1

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rather stoichiometric composition without variation of the lattice parameters. These findings were confirmed by [1993Bau, 1996Hil2]. Details of the crystal structure with linear C-B-C chains are given by [1996Hil2]. Lattice parameters of Al3BC3 were measured at room temperature up to 7.5 GPa using a multi-anvil synchrotron system with B4C anvils; for a high temperature pressure experiment the sample was placed in a graphite ampoule [2000Sol]. Al3BC3 is free of structural transitions up to 1523˚C within the pressure range 2.5 to 5.3 GPa [2000Sol]. A further ternary compound τ5 was observed after infiltration by liquid Al at 1170˚C with post heat treatment for 100 h at 800 to 1000˚C [1987Sar]. The hexagonal lattice was established by TEM; the approximate composition “Al4BC” resulted from EELS-data [1987Sar]. This phase has been also confirmed by [1989Hal, 1990Pyz]. From a detailed investigation of this Al-rich boroncarbide by X-ray powder diffraction, LOM and EPMA, [1992Via] suggested a formula of Al3BC rather than “Al4BC” and attributed a hexagonal cell; additional weak lines in the X-ray intensity pattern of Al3BC prompted a larger unit cell pffiffiffi a = a0/ 3 [1993Gon]. Although the authors of [1997Mey] recognized the larger cell, the crystal structure of Al3BC was solved for the high symmetry subcell from single crystals isolated from a sample directly reacted from the elements - however, from EPMA a composition Al2.5BC was derived (see also Table 2). Al6B-octahedra and trigonal Al5C-bipyramids are the characteristic structural elements [1997Mey]. The various data on the compositional ranges of the τ4 and τ5 phases are summarized in Fig. 2. Half filled circles correspond to the accepted stoichiometries Al3BC and Al3BC3. From the significant change of the unit cell volume of Al4C3 comparing binary and ternary alloys, a solubility of boron is suggested [1996Bid, 2002Zhe, 2000Mey]. Solubility of boron in Al4C3 was established to be 3.4 at.% at 900˚C [2002Zhe] and an interesting behavior of lattice parameters was observed. In spite of the increase of the “a” parameter and of the cell volume with boron content, the “c” parameter decreases. That may be explained by a preferential distribution of boron and carbon atoms among different crystallographic sites. A significant solubility of boron in Al4C3 was also reported by [2000Mey] to be about 9.3 at.%, however, no details on the relevant temperature were given. Furthermore these authors claim for Al4C3–xBx lattice parameters increasing with boron content. Lattice parameters of Al4C3 for samples located in three-phase regions (Al)+Al4C3+Al3BC, Al4C3+Al3BC3+Al3BC and Al4C3+Al3BC3+(C) are very close, assuming that these three-phase regions meet at the Al4C3 phase at a maximal boron solubility of Al4(C0.92B0.08)3. Insignificant solubility of carbon in AlB2 is reported by [2002Zhe] comparing lattice parameters in ternary and binary samples; AlB2 with 0.5 at.% C, heat treated at 900˚C, already contains the Al3BC phase.

Isothermal Sections Phase equilibria for the 1400˚C isothermal section are summarized in Fig. 3, revealing the existence of four ternary compounds τ1 to τ4. A small field of liquid phase exists at 1400˚C which is in equilibrium with Al3B48C2, Al2.1B51C8 and with Al3BC3 [1993Bau]. Boron-poor equilibria agree with an earlier work by [1980Ino] who reported on the two-phase equilibria Al4C3+Al3BC3 (Al7B4C8), Al3BC3 (Al7B4C8)+B4C and Al3BC3 (Al7B4C8)+C. In Fig. 3 twophase equilibria are shown to exist between the binary solid solution “B4C” and AlB40C and Al2.1B51C8. At 1400˚C all ternary compounds seem to exist at their stoichiometric compositions [1993Bau], whilst [1965Eco] claimed a homogeneity range for τ1 at 1800˚C from DOI: 10.1007/978-3-540-88053-0_3 ß Springer 2009

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“AlB48C8” to Al3B48C8. Binary AlB12 was never seen in combination with Al2.1B51C8 nor with AlB40C4 [1993Bau]. The isothermal section at 1000˚C, Fig. 4, was constructed on the basis of data from [1997Via]. Due to low interaction kinetics in the boron- and carbon-rich part of the system at 1000˚C, equilibria in this portions of the diagram are preliminary. Moreover, ternary compounds τ1 and τ2 were not included in the 1000˚C section by [1997Via], τ4 was listed as “Al8B4C7”, and no solubility of boron in Al4C3 was considered. For consistency with the present knowledge on the Al-B-C system, the ternary compounds τ1 and τ2 were introduced in Fig. 4 and the composition of τ4 was changed to Al3BC3. The solubility of boron in Al4C3 at 1000˚C was estimated to be about 4 at.%, extrapolating from data of [2002Zhe] at 900˚C. Figure 5 represents the isothermal section at 900˚C [2002Zhe] confirming the equilibrium AlB2+Al3BC, whereas [1997Via] claimed this equilibrium to be only stable below 868 ± 4˚C. Similar to 1000˚C the equilibria at 900˚C involving τ1 and τ2 are not well established due to low reactivity of the components.

Invariant Equilibria, Liquidus Surface A tentative liquidus surface for the aluminum rich portion of the diagram (Fig. 6) was proposed by [1997Via], presenting equilibria involving the τ5 phase. The invariant equilibrium U5 (L+Al3B48C2ÐAlB2+Al3BC) was reported at 868 ± 4˚C by [1997Via], but this temperature can not be accepted in respect to the observed isothermal equilibrium AlB2+Al3BC at 900˚C [2002Zhe] suggesting such transformation above 900˚C. Comparison of the reaction scheme and the isothermal section at 1000˚C (Fig. 4) with the isothermal section at 1400˚C (Fig. 3) suggests a rather complicate picture of the phase transformations in this regions mainly due to decomposition of τ5. Based on an earlier thermodynamic calculation by [1982Doe], a reaction scheme was derived [1990Luk], which gives a tentative information of the solidification behavior in the AlB-C ternary. The temperatures of the invariant equilibria were estimated and the ternary compounds τ1 to τ3 were assumed to be part of the solid solution arising from binary B4C; τ5 was not included. A more recent thermodynamic modelling of the Al-B-C phase diagram by [1993Wen] as part of the multi-component Al-B-C-N-Si-Ti system treated the ternary compounds τ1, τ2, τ3 as independent phases, however, the peritectoid formation of τ4 (Al3BC3) is in strict contradiction to the experimentally confirmed two-phase equilibrium τ4+τ5 (Al3BC3+Al3BC) [1997Via, 2002Zhe] as well as to the observed existence of τ4+τ5 in as cast alloys [2002Zhe], thereby strongly indicating direct formation of Al3BC3 from the liquid. A closed ternary miscibility gap in the Al-rich liquid is suggested from thermodynamic calculations by [1993Kau], however, hitherto without experimental confirmation [2002Zhe]. Figure 7 presents a reaction scheme for the major parts of the Al-B-C phase diagram. The reaction scheme is essentially based (i) on the tentative liquidus projection for the Al-rich part as suggested by [1997Via], (ii) on the experiments of [2002Zhe] concerning the solidification of the phases τ4, τ5 and (iii) on the thermodynamic calculation of [1993Wen] for the B-rich part, however, accepting peritectic formation of AlB12.


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Thermodynamics Enthalpies of formation and heat capacity measurements from a Calvet type automatic microcalorimeter in the temperature range 310–1200 K were reported by [1987Kis] and are listed as follows: Al3 B48 C2 : H 0 ðTÞ H 0 ð298Þ ¼ 0:7945 103 T 2 þ 0:5182T 225:1374 ðin J g1 Þ and Cp ðTÞ ¼ 0:1589 102 T þ 0:5182 ðJ g1 K1 Þ Al2:1 B51 C8 : H 0 ðTÞ H 0 ð298Þ ¼ 0:7226 103 T 2 þ 0:5411T 225:5702 1

ðin J g

:

for AlB24 C4 Þ and

Cp ðTÞ ¼ 0:1589 102 T þ 0:5182 ðJ g1 K1 for AlB24 C4 Þ Thermodynamic calculations of the Al-B-C system due to [1982Doe, 1993Wen, 1993Kau], however, are not fully consistent with experimental observations. For detailed discussion, see section Invariant Equilibria.

Notes on Materials Properties and Applications Mechanical properties of Al-B4C cermets and boron/carbon fiber-aluminium composites have been investigated by various groups [1972Bak, 1973Her, 1975Mun, 1984Via, 1985Che, 1985Hal, 1985Kov, 1985Pyz, 1985Sar, 1986Che, 1986Dub,1990Ram, 1996Pyz, 2002Ars]; the effect of reaction on the tensile behavior of infiltrated composites was reported by [2002Kou2] and size dependent strengthening in particle reinforced Al by [2002Kou1]; reaction products were studied by [2001Lee]. An increase of surface hardness of about 25 to 40 % can be achieved by impulse laser radiation on B4C/Al cermets [1988Kov]. Wetting of B4C by Al has been studied by many research teams with rather contradicting results, until the temperature and time dependent occurrence of chemical reactions/compound formation was analyzed in detail (for discussion see i.e. [1979Kis, 1979Pan, 1989Hal, 2000Kha]). The kinetics of wetting by liquid aluminium of flat, sintered boron carbide specimens with residual porosity less than 3 % were investigated by [1979Pan]. The speed of spreading of liquid aluminium at 1100˚ to 1200˚C was measured to be 0.1–0.8 mm·s–1, in accordance with r2 = f(t), where r equals the radius of the contact circle. The angle of contact was first 92˚, however, in 3 to 5 min decreased to 28˚. The slow spreading was determined by the formation of new aluminum boron carbide phases in the contact zone with a microhardness of ca. 13 GPa. The driving force Δσ = σ (cos Θ0-cosΘ) (σ = surface tension of the melt, Θ0 = contact angle of the melt, Θ = contact angle at time (t)) decreased sharply becoming zero in 4 to 5 min [1979Pan]. The contact angle of molten Al on B4C as a function of processing time for various isotherms at 5·10–3 to 10–4 Pa was also given by [1989Hal] based on sessile drops cooled to room temperature. Mechanical properties, electrical and thermal conductivity as well as their temperature dependencies were reported on the Knoop and Vickers microhardness for Al-flux grown (temperature region 1750 to 800˚C) ‘‘amber’’ single crystals Al3B48C2 and for ‘‘black’’ crystals (αAlB12, γAlB12 and AlB2.1B51C8) [1986Kis]. These studies were also performed on hot pressed specimens of various compositions x(AlB12)+(1–x)B4C and Al3B48C3 in the temperature range 24 to 827˚C [1991Kha1, 1991Kha3]. For Al3BC3 (“Al8B4C7”), Al3B48C3, Al2.1B51C8 [1991Kha2] DOI: 10.1007/978-3-540-88053-0_3 ß Springer 2009

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also examined these properties as a function of porosity and quantity of Fe-impurity. These data are summarized in Table 3 including information on flux-grown crystals Al3B48C2 and Al2.1B51C8 [1990Oka]. Both types of crystals were said to be p-type semiconductors [1986Kis]. In a ring test the strength of a powder compact of B13C2 +1 mass% Al, sintered at 1950˚C, was found to be 0.50(7) GNm–2 [1978Ekb]. [1991Kha2] reported on the kinetics of thermal densification of hot pressed powders of B4C, AlB12, Al3B48C2 and Al3BC3. Kinetics of dissolution in HCl, HNO3 and HCl-HNO3 was studied by [1998Kha] as well as the resistance of Al-boron carbides to alkali and hydrogen peroxide. [1989Hal] studied the densification kinetics of Al+B4C cermets in the range from 800 to 1400˚C in pressureless sintering as well as after applying hot isostatic pressure. The kinetic of metal depletion in post heated dense cermets B4C/Al at temperatures between 600˚C and 1000˚C was investigated by [1990Pyz]. Chemical stability against various boiling acids, oxidation resistance, IR and EPR spectra of Al-borides and Al-boron carbides (Al3B48C2, Al2.1B51C8) was studied by [1991Pri]. The spectra were taken at 77 K and 300K and for different crystal orientations relative to the magnetic field. Absorption edge and IR-active phonons in Al3B48C2 were reported by [1987Hau, 2000Wer] and IR spectra of boron carbide containing up to 1.5 at.% Al were determined between of 8 to 500 mm–1 wave numbers and for temperatures between 70 to 450 K [1997Sch]. These data seem to suggest the incorporation of Al-atoms into binary boroncarbide in form of pairs substituting the B-B-C or C-B-C chains [1997Sch]. Characteristic IR absorption bands for finely dispersed powders of Al-borides and Al-boron carbides were listed by [1998Kha]. The Seebeck-coefficients were reported to linearly increase from 260 μVK–1 for binary “B4C” to 450 μVK–1 for 1.4 at.% Al dissolved, revealing p type behavior [1997Sch]. Seebeckcoefficients, thermal and electric conductivities were further reported by [2000Liu] for B4.3Cbased samples containing 0.5, 10, 15, 20 mass% Al, highlighting the Z-value at RT of 1.04·10–6K–1 for the 5 mass% Al sample. IR and Raman spectroscopy on Al3BC3 (at RT) confirm the linear (CBC)5– unit as an isoelectronic CO2-analogon [1996Hil2, 2000Mey]. On heating in air, Al3BC3 (earlier reported as Al8B4C7), Al3B48C3 and Al2.1B51C8 show low oxidation at 500˚C (increase of mass 4 mgh–1); intensive oxidation, with a mass increase of 40 mgh–1) starts at 1280˚C for Al2.1B51C8 and at 1370˚C for Al3B48C2 [1991Pri, 1989Kha, 1991Kha4]. Oxidation in air of single crystals Al2.1B51C8 and Al3B48C2 started at about 760˚C and 710˚C, respectively [1990Oka]. The reaction products were 9Al2O3·2B2O3 for Al2.1B51C8 crystals and B2O3 for AlB40C4 specimens [1994Kud]. Whereas Al3BC3 was said to be unstable in acids [1991Kha4], more detailed experiments [1996Hil2] proved stability at room temperature against bases and dilute acids, except for HNO3 and HF. Al3BC3 was furthermore said to be stable in air up to 600˚C [1991Kha4,1996Hil2]. Al3BC is quickly attacked by dilute HCl [1997Mey]. Thermophysical properties of sintered bodies of Al3BC3 have been derived by [2000Wan]. These are linear thermal expansion in the range of 25 to 1200˚C, specific heat and thermal diffusivity via laser flash technique, Youngs modulus of 136.6 GPa, Vickers hardness of 12.1 GPa at a load of 196 N and thermogravimetric recording of growth of an oxidized layer on heating in air up to 1500˚C. Fitting a Birch-Murnaghan equation of state to the pressure dependency of the lattice parameters of Al3BC3 up to 7.5 GPa, the isothermal bulk modulus was B0 = 153 ± 6 GPa (dB0/ dp = 19 ± 4) [2000Sol]. Despite high bulk modulus the Vickers hardness of single crystals is as low as 20.7 GPa at a load of 25g and 18.2 GPa at a load of 50g [2000Sol]. Landolt‐Bo¨rnstein New Series IV/11E1

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Al3BC was successfully prepared by self propagating high temperature synthesis induced by mechanical activation of Al-B-C powder mixtures in air; mixtures low in boron (AlB0.1C) resulted in Al3BC3 under violent emission of heat [1999Tsu]. In contrast to that [2000Sav] was unable to prepare ternary aluminoborocarbides from mechanochemical synthesis. Elastic bulk and shear moduli for Al3BC (earlier reported as Al4BC) were measured by [1995Pyz] and estimated by [1999Tor].

Miscellaneous A series of patents covers the techniques to produce dense B4C/Al cermets by infiltration of the metal matrix into the porous ceramic body without wetting reactions [1976Lan, 1986Hal, 1987Pyz, 1990Pyz, 1991Pyz, 1995Pyz, 1996Pyz, 1997Du , 2000Pyz, 2001Lee]; subsequent heat treatment results in materials with designed chemistry and microstructures, flexure strength, hardness and fracture toughness. Fine microstructures were obtained via ultrarapid microwave heating [1995Rug]. B4C/Al cermets have been considered as an improved structural neutron absorber [1977Ros, 1978Boi, 1978Sur, 1986Ros, 1987Lev, 1992Bei] and for applications as friction materials for automotive brake applications [1999Cha]. Oxidation protective B4C-coatings on C-fibers in Al-matrix were reported by [1996RMi] and [1996Vin] produced C-fibres-Al composites by a squeeze casting technique. Explosive consolidation to produce Al/ B4C composites was studied by [1995Bon, 1997Yue]. Shock recovery experiments were performed on a 65 vol% B4C-Al cermet as a function of shock pressure [1989Blu].

. Table 1 Literature Data on Experimental Temperatures of Invariant Equilibrium L+AlB12ÐAlB2 T [˚C]

Technique

Heating Rate

References

Stability observation

-

1000 - 1500

[1936Hof]

Synthesis observation

-

980

[1967Ser]

DTA

4˚C/min

920

[1967Ato]

Stability observation

-

1350 - 1500

[1972Sir]

DTA

5˚C/min

1030 ± 5

[1993Ips]

Synthesis observation

-

892 ± 5

[1997Via]

DSC and Stability observation

0˚C/min*

956 ± 5

[2000Hal]

DSC

10˚C/min

914 ± 55

[2001Fje]

* DSC measurements were performed with heating rate of 5, 15 and 40˚/min., and extrapolated to 0˚C/min.

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. Table 2 Crystallographic Data of Solid Phases

Phase/ Temperature Range [˚C]

Pearson Symbol/ Space Group/ Prototype

Lattice Parameters [pm]

Comments/References

(Al) < 660.452

cF4 Fm3m Cu

a = 404.96

[Mas2 ]

(βB) < 2092

hR333 R3m βB

a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1096.5 c = 2386.8 a = 1097.4 c = 2387.7

[Mas2 , 1993Wer]

a = 246.12 c = 670.90 a = 246.023 c = 671.163 a = 246.75 c = 669.78

[Mas2 ]

(C) < 3827 (B.P.)

B4C < 2450

hP4 P63/mmc C-graphite

hR45 R3m B13C2

at 1.1 at.% C [1993Wer] linear da/dx, dc/dx at AlB31 [V-C2 ] from sample Al4B95C1, quenched from 1400˚C, contains Al3B48C2 and αAlB12 [1993Bau]

[1967Low] at 2.35 at.% Bmax (2350˚C) linear da/dx, dc/dx, [1967Low]

a = 565.1 c = 1219.6 a = 560.7 c = 1209.5 a = 560.3 c = 1209.8

from sample containing τ2, τ4, quenched from 1400˚C [1993Bau]

9 to 20 at.% C [1990Ase]

B25C

tP52 P42m B25C

a = 872.2 c = 508.0

[V-C2 ] also B51C1, B49C3; all metastable?

Al2B3 ≤ 525

hR* Al2B3 (?)

a = 1840 c = 896

at 60 at.% B [1992Var] metastable?

AlB2 ≤ 956±5

hP3 P6/mmm AlB2

a = 300.6 b = 325.2 a = 300.67 ± 0.01 b = 325.36 ± 0.02 a = 300.63 ± 0.01 b = 325.46 ± 0.01 a = 300.43 ± 0.03 b = 325.19 ± 0.06

[1994Dus], temperature from [2000Hal] [2002Zhe]


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. Table 2 (continued)

Phase/ Temperature Range [˚C] αAlB12 ≤ 2050

γAlB12

Pearson Symbol/ Space Group/ Prototype tP216 P41212 αAlB12

oP384 P212121 γAlB12

Lattice Parameters [pm] a = 1015.8 c = 1427.0 a = 1018 c = 1434.3 a = 1016.3 c = 1425.6 a = 1015.5 c = 1426.0 a = 1014.93 ± 0.07 c = 1425.0 ± 0.5 a = 1014.4 b = 1657.3 c = 1751.0

hR21 R 3m Al4C3

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[1994Dus] ρexp. = 2.65 Mgm–3 [1991Pri] from sample Al2B92C2, quenched from 1400˚C, contains Al3B48C2 [1993Bau] from sample Al4B95C1, quenched from 1400˚C, contains Al3B48C2 and AlB31 [1993Bau] [2002Zhe] [1983Hig, 1994Dus, 2000Hig] metastable phase or ternary product stabilized by small amounts of impurity metals present in Alflux grown material ρexp. = 2.56 Mgm–3 [1991Pri]

a = 1019.5 b = 1666 c = 1769 Al4C3 < 2156

Comments/References

a = 333.8 c = 2511.7 a = 334.21 ± 0.01 c = 2503.2 ± 0.5 a = 335.78 ± 0.02 c = 2499.6 ± 0.5

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Phase/ Temperature Range [˚C] * τ1, Al2.1B51C8 (eventually low temperature phase of τ2)

* τ2, AlB40C4 (eventually high temperature phase of τ1)

Pearson Symbol/ Space Group/ Prototype oC88 Cmcm Al2B51C8

hR45 R3m B4C-deriv.


earlier labeled “AlB10” [1967Wil] or AlB24C4 [1964Mat, 1969Wil, 1970Wil] [1969Per] ρexp. = 2.54 Mgm–3

a = 569.0 b = 888.1 c = 910.0 a = 568.7 b = 887.7 c = 909.8 a = 569.0 b = 888.1 c = 910.0 a = 569.3 b = 884.7 c = 909.3 a = 567.6 b = 891.4 c = 909.5 a = 569.2 b = 889.2 c = 911.2

from sample containing τ2 and τ3, quenched from 1400˚C [1993Bau] from sample containing τ4, quenched from 1400˚C [1993Bau] from sample Al4B92C4 quenched from 1400˚C, contains Al3B48C2 (tetragonal), Al3B48C2 (A) and AlB40C4 [1993Bau] [1991Pri]

[1990Oka] ρexp. = 2.54 Mgm–3 single crystals from Al-flux

a = 564.2 c = 1236.7 a = 565.37 c = 1231.4 a = 564.8 c = 1239.9 a = 565.6 c = 1238.9

[1970Nei] ρexp. = 2.52 Mgm–3 from sample containing τ1, τ3, quenched from 1400˚C [1993Bau] from sample containing τ4 and B4C, quenched from 1400˚C [1993Bau] from sample Al4B92C4 quenched from 1400˚C, contains also Al3B48C2 (tetrag.), Al3B48C2 and Al2.1B51C8 [1993Bau] [1966Gie] for composition “Al2B48C8”

a = 563 c = 1129 a = 565 c = 1239


Comments/References

[1966Lip] for composition “Al4B48C8”

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Phase/ Temperature Range [˚C] * τ3, Al3B48C2 (r) < 650

Pearson Symbol/ Space Group/ Prototype oI212 Imma Al3B48C2

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Lattice Parameters [pm] a0 = 1240.7 b0 = 1262.3 c0 = 1014.4 a = 1234 b = 1263 c = 508 a = 1232.5 b = 1261.4 c = 1016.2 a = 1233.72 b = 1262.41 c = 1016.06 a = 1232.5 b = 1264.7 c = 1016.2 a = 1230.2 b = 1262.1 c = 1016.1 a = 1229.1 b = 1262.2 c = 1015.88 a = 1233.62 b = 1262.40 c = 1015.94 a = 1239.0 ± 0.3 b = 1263.7 ± 0.3 c = 1013.6 ± 0.4 a = 1237.7 b = 1262.7 c = 507.9 to a = 1236.3 b = 1261.6 c = 510.2 a = 616.6 b = 1263.5 c = 1065.6 a = 618.1 b = 1262.2 c = 1016.1 a = 617 b = 1263 c = 1016

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Comments/References [1996Hil1], only one low temperature modification! [1965Mat], two modifications, microscopically twinned; modification A, c = c0/2 from a sample Al6B92C2 cooled from 1400˚C contains “AlB12” [1993Bau] from sample Al4B95C1 cooled from 1400˚C contains also “AlB12”, AlB31 [1993Bau] [1991Pri]

[1994Kud]

from sample Al4B92C4 cooled from 1400˚C, [1993Bau] contains Al2.1B51C8, AlB40C4 and tetragonal Al3B48C2 [1993Bau] from sample Al4B95C1, cooled from 1400˚C, see above. [2000Wer]

[1990Oka] single crystals from Al-flux modification A ; c = c0/2

[1990Oka] single crystals from Al-flux ρexp. = 2.59(2) Mgm–3 modification B, a = a0/2

[1965Mat] modification B a = a0/2


Al–B–C

3




Lattice Parameters [pm] a = 616.4 b = 1262.1 c = 1016.4

* τ3, Al3B48C2 (h) > 650

* τ4, Al3BC3 < 1835

*τ5, Al3BC < 1100

[1991Pri] a = a0/2

tP52 a = 885 P42/nnm c = 508 B25C -deriv. a = 882 c = 509 a = 881.9 c = 508.25 hP42 P3c1a) Mg3BN3

hP20 P3c1 (P63/mmc for subcell) Al3BC

Comments/References

[1996Hil1] high temperature modification [1965Mat] from sample Al4B92C4 cooled from 1400˚C, contains also Al2.1B51C8, AlB40C4 and orthorhombic Al3B48C2 [1993Bau]

a = 589.97 c = 1589.0 a = 590.6 c = 1590.1 a = 590.7 c = 1591.3 a = 590.5 c = 1590.5 a = 340.1 ± 0.3 c = 1584 ± 0.2 a = 590.22 ± 0.3 c = 1589.4 ± 0.1 a = 605.0 c = 1154.0 a = 603.45

[1996Hil2] ρ = 2.66 Mgm–3 temperature from [1980Ino] [1980Ino], labelled as Al8B4C7 from sample containing τ1, quenched from 1400˚C [1993Bau] from sample containing τ2 and B4C, quenched from 1400˚C [1993Bau] pffiffiffi [2000Sol], subcell with a = a0/ 3 pressure dependence of the lattice parameters is given up to 7.5GPa [2002Zhe] [1993Gon, 1997Via]

c = 1152.02 a = 6041.9 ± 0.2 c = 1154.0 ± 0.3 a = 349.1 c = 1154.1 a = 352.0 c = 582.0

[1997Mey] from single crystals, “Al2.5BC” from EPMA [2002Zhe] pffiffiffi [1992Via], subcell with a = a0/ 3 [1987Sar], earlier “Al4BC” subcell with a = a0/ p ffiffiffi 3, c = c0/2

pffiffiffi P63/mmc for subcell with a = a0/ 3, c = c0

a)


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(101)

(001)

γAlB12

(100)

(111)

(100)

(111)

αAlB12

Al3BC3

Al2.1B51C8

Al3B48C2

29.6(1.0) 25.8(7) 22.8(8) 21.6(1.1)

19.6(5)

22.6 26 6 24.2(7) 25.0–26.9

37.6(2.0) 31.7(8) 23.7(6)

27.1(5)

26.5(5) 23.1

(0.5N, 293K) (1N, 293K) (2N, 293K) (4.9N, 293K)

(2N, 293K)

(2N, 293K) (5N, 293K) (5N, 1200K) (2N, 293K) (1N, 293K)

(0.5N, 293K) (1N, 293K) (4.9N, 293K)

(2N, 293K)

(2N 293K) (5N, 293K)

34.4(2.7) 31.0(1.5) 27.3(1.2) 23.8(9)

20.7 18.2

25.7 to 30.5

(0.5N, 293K) (1N, 293K) (2N, 293K) (4.9N, 293K)

(0.25N, 293K) (0.50N, 293K)

(1N, 293K)

33.6(1.6) (2N, 293K)

1.8(2)

1.5(3)

2.7(2)

5.3 4(1)

Fracture Toughness K1c [MPa·m1/2]

0.1

1 0.6 - 1.2

0.18 - 0.36

3.85 · 105 0.22

5.92 ·102

2.02 · 105 0.08 - 0.18

10–3 - 1

104-106 2.6·10310–6 2.6·10–6

ρ293K [Ωm]

Activation Energy ΔE [eV] 100K to 400K R=R0exp(-ΔE/2kT)

38.7 (310K) 60 (600K)

19.6 (310K)

Thermal Conductivity [Wm–1K–1]

[1986Dub] [1986Dub] [1991Pri] [1986Dub]

[1991Pri]

[2000Sol]

[1986Kis] [1986Kis] [1986Kis] [1991Pri] [1990Oka]

[1990Oka, 1994Kud] [1986Dub] [1986Dub] [1986Dub]

[1986Kis] [1986Kis] [1986Kis] [1991Pri]

References

3

Microhardness [GPa] at Various Loads and Temperatures Crystal Face Vickers Compound of Indent Knoop

. Table 3 Microhardness, Fracture Toughness, Electrical Conductivity, Activation Energies for Electrical Conductivity and Thermal Conductivity for Various Aluminium Borides and Aluminium Boron Carbides

14 Al–B–C


Al–B–C

3

. Fig. 1a Al-B-C. Various versions of the Al-rich part of the Al-B diagram


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. Fig. 1b Al-B-C. Accepted Al-B phase diagram

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Al–B–C

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. Fig. 1c Al-B-C. Accepted Al-B phase diagram, enlarged Al-rich region


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Al–B–C

. Fig. 2 Al-B-C. Superposition of literature data on the homogeneity regions of τ4 and τ5 phases. Halffilled circles correspond to the accepted compositions

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Al–B–C

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. Fig. 3 Al-B-C. Isothermal section at 1400˚C; the position of Al3BC is indicated by a full circle


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. Fig. 4 Al-B-C. Isothermal section at 1000˚C

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Al–B–C

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. Fig. 5 Al-B-C. Isothermal section at 900˚C


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. Fig. 6 Al-B-C. Tentative liquidus surface projection

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. Fig. 7 Al-B-C. Reaction scheme

Al–B–C


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References [1936Hof] [1960Koh] [1963Wil] [1964Mat] [1965Eco] [1965Mat] [1966Gie] [1966Lip] [1967Ato] [1967Bli] [1967Low] [1967Ser] [1967Wil] [1969Per] [1969Wil] [1970Nei] [1970Wil] [1972Bak] [1972Sir] [1973Her]

[1975Mun]

[1976Lan]

[1976Mon] [1977Mat] [1977Ros]

Hofmann, W., Jaeniche, W., “Contribution to the Knowledge of the Aluminium-Boron System” (in German), Z. Metallkd., 1, 1–5 (1936) (Phase Diagram, Crys. Structure, 13) Kohn, J.A., Eckart, D.W., “Aluminium Boride, AlB12”, Anal. Chem., 32, 296–298 (1960) (Crys. Structure, Experimental, 6) Will, G., “On the Crystal Structure of AlB10”, J. Am. Chem. Soc., 85, 2335–2336 (1963) (Crys. Structure, Experimental; 7) Matkovich, V.I., Economy, J., Giese Jr. R.F., “Presence of Carbon in Aluminium Borides”, J. Am. Chem. Soc., 86, 2337–2340 (1964) (Crys. Structure, Experimental, 14) Economy, J., Matkovich, V.I., Giese, Jr.R.F., “Crystal Chemistry of α-Boron Derivatives”, Z. Kristallogr., 122, 248–258 (1965) (Review, Crys. Structure, 26) Matkovich, V.I., Giese, Jr.R.F., Economy, J., “Phases and Twinning in C2Al3B48”, Z. Kristallogr., 122, 108–155 (1965) (Crys. Structure, Experimental, 7) Giese, Jr.R.F., Economy, J., Matkovich, V.I., “Topotactic Transition in C4AlB24”, Acta Crystallogr., 20, 697–698 (1966) (Crys. Structure, Experimental, 7) Lipp, A., Ro¨der, M., “On an Aluminium Bearing Boron Carbide” (in German), Z. Anorg. All. Chem., 343, 1–5 (1966) (Crys. Structure, Experimental,13) Atoda, T., Higashi, I., Kobayashi, M., “Process of Formation and Decomposition of Aluminium Borides”, Sci. Papers Inst. Phys. Chem. Res., 61, 92–99 (1967) (Phase Diagram, Crys. Structure, 8) Bliznakov, G., Peshev P., Niemyski, T., “On the Preparation of Crystalline Aluminium Borides by a Vapour Deposition Process”, J. Less-Common Met., 12, 405–410 (1967) (Experimental, 14) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Serebryanskii, V.T., Epel’baum, V.Z., Zhdanov, G.S., “Equilibrium Diagram of the Aluminium - Boron System”, Russ. J. Inorg. Chem., 12(9), 1311–1316 (1967) (Phase Diagram, 33) Will, G., “Crystal Structure Analysis of AlB10 by the Convolution Molecule Method”, Acta Crystallogr., 23, 1071–1079 (1967) (Crys. Structure, 11) Perrotta, A.J., Townes, W.D., Potenza, J.A., “Crystal Structure of C8Al2.1B51”, Acta Crystallogr., 25B, 1223–1229 (1969) (Crys. Structure, Experimental, 11) Will, G., “The Crystal Structure of C4AlB24”, Acta Crystallogr., 25B, 1219–1222 (1969) (Crys. Structure, Experimental, 11) Neidhard, H., Mattes, R., Becher, H.J., “On the Preparation and Structure of an Aluminium Bearing Boron Carbide”, Acta Crystallogr., 26B, 315–317 (1970) (Crys. Structure, Experimental, 11) Will, G., “On the Existence of AlB10: a Critical Review of the Crystal Structures of AlB10 and C4AlB24”; Electrochem. Technol., 3(1-2), 119–126 (1970) (Crys. Structure, Experimental, 11) Baker, A.A., Braddick, D.M. Jackson, P.W., “Fatigue of Boron-Aluminium and Carbon-Aluminium Fibre Composites”, J. Mater. Sci., 7, 747–62 (1972) (Mechan. Prop., Experimental, 18) Sirtl, E., Woerner, L.M., “Preparation and Properties of Aluminium Diboride Single Crystals”, J. Cryst. Growth, 16, 215–218 (1972) (Crys. Structure, Phase Diagram, 15) Herring, H.W., Lytton, J.L., Steele, J.H., “Experimental Observations of Tensile Fracture in Unidirectional Boron Filament Reinforced Aluminium Sheet”, Metall. Trans. A, 4(3), 807–817 (1973) (Experimental, Mechan. Prop., 9) Munir, Z.A., Veerkamp, G.R., “Investigation of the Parameters Influencing the Microstructure of HotPressed Boron Carbide”, California Univ., Davis (USA). Dept. of Engineering,. 95 pp. (1975) (Mechan. Prop., Crys. Structure, 32) Lange, R.G., Munir, Z.A., “Sintering Kinetics of Pure and Doped Boron Carbide. Final Technical Report”, California Univ., Davis (USA). Dept. of Mechanical Engineering, 35 pp. (1976) (Experimental, 0) Mondolfo, L.F., “Aluminium - Boron System” in “Aluminium Alloys: Structure and Properties”, Butterworths, London, pp. 228–230 (1976) (Review, Phase Diagram, 29) Matkovich, V.I., Economy, J., “Structural Determinants in Higher Borides” in “Boron and Refractory Borides”, Matkovich, V.I. (Ed.), Springer Verlag, Berlin, 78–95 (1977) (Crys. Structure, Review, 36) Roszler, J.J., “Production of Neutron Shielding Material. Patent; B4C+Al in Al Boxes”, US Patent Document 4,027,377/A/, (1977)

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Al–B–C [1978Boi]

[1978Ekb] [1978Sur] [1979Kis]

[1979Pan]

[1980Ino] [1982Doe]

[1983Hig] [1984Sig] [1984Via] [1985Che]

[1985Hal]

[1985Kov]

[1985Pyz]

[1985Sar]

[1986Che] [1986Dub] [1986Hal] [1986Kis] [1986Pes]

[1986Ros] [1987Hau]

3

Boiko, Yu.V., Gol’tsev, V.P., Gorobtsov, V.G., Kavkhuta, G.A., Strelkov, G.I., Khrenov, O.V., Yuzhanin, M.I., “Development and Investigation of Properties of Disperse Boron-Containing Materials for Control Rods of a Nuclear Reactor” (in Russian), Vest. Akad. Navuk BSSR, Ser. Fiz.-Energ. Navuk, 3, 5–8 (1978) (Mechan. Prop., Experimental) Ekbom, L.B., “Effect of Increased Boron Content on the Sintering Behavior and Mechanical Properties of Boron Carbide”, Keram. Z., 183–189 (1978) (Experimental, Mechan. Prop., 6) Suri, A.K., Gupta, C.K., “Studies on the Fabrication of Aluminium Bonded Boron Carbide Rings”, J. Nucl. Mater., 74(2), 297–302 (1978) (Experimental, 4) Kislyi, P.S., Kozina, G.K., Bodnaruk, N.I., “Wetting and Impregnation of Boron Carbide with Copper, Aluminum, and Their Alloys” (in Russian), Adgez. Rasplav. Pajka Mater., 4, 54–57 (1979) (Experimental) Panasyuk, A.D., Oreshkin, V.D., Maslennikova, V.R., “Study of the Kinetics of the Reactions of Boron Carbide with Liquid Aluminium, Silicon, Nickel and Iron”, Sov. Powder Metall. Met. Ceram., 199(7), 487–490 (1979), translated from Poroshk. Metall., 199(7), 79–83 (1979) (Experimental, 9) Inoue, Z., Tanaka, H., Inomata, Y., “Synthesis and X-Ray Crystallography of Aluminium Boron Carbide”, J. Mater. Sci., 15, 3036–3040 (1980) (Crys. Structure, Experimental, 7) Do¨rner, P., “Constitutional Investigations on High Temperature Ceramics of the B-Al-C-Si-N-O System by Means of Thermochemical Calculations” (in German), Thesis, Univ. Stuttgart (1982) (Experimental, Thermodyn., 126) Higashi, I., “Aluminum Distribution in the Boron Framework of γ-AlB12”, J. Solid State Chem., 47, 333–349 (1983) (Crys. Structure, 17) Sigworth, G.K., “The Grain Refining of Aluminium and Phase Relationships in the Al-Ti-B System”, Mater. Trans. 15A, 277–282 (1984) (Experimental, Phase Diagram, Thermodyn. Calculation, 28) Viala, J. C., Bouix, J., “Elaboration of Aluminum-Matrix Composite Materials Reinforced with Inorganic Fibers”, Mater. Chem. Phys., 11(2), 101–123 (1984) (Mechan. Prop., Experimental, 41) Chernyshova, T.A., Tsirlin, A.M., Gevlich, S.O., Rebrov, A.V., Obolenskii, A.V., “Effect of Surface Condition on the Strength of Coated Boron Fibers”, Sov. Powder Metall. Met. Ceram., 24(3), 210–213 (1985), translated from Poroshk. Metall., 24(3), 39–43 (1985) (Mechan. Prop., Experimental, 9) Halverson, D.C., Pyzik, A.J., I.A. Aksay, I.A., “Processing and Microstructural Characterization of B4CA1 Cermets”, “Composites and Advanced Ceramic Materials”, Anon. Proc. 9th Annu. Conf., American Ceramic Society, Inc., Columbus, OH, 736–744 (1985) (Mechan. Prop., Experimental, 14) Koval’chenko, M.S., Laptev, A.V., Zhidkov, A.B., “Annealing Effect on Structure and Properties of Hot Pressed Cermets Based on Boron Carbide” (in Russian), Poroshk. Metall., 24(9), 51–54 (1985) (Mechan. Prop., Experimental, 6) Pyzik, A. J., Aksay, I. A., “Processing, Microstructure, and Mechanical Properties of Boron CarbideAluminum Alloys Composites”, Anon. Abst. 38th Annu. Pacific Coast Regional Meeting American Ceramic Society, American Ceramic Society, Columbus, OH, (1985) (Mechan. Prop., Experimental, 0) Sarikaya, M., Pyzik, A.J., Ilsay, I.A., Snowden, W. E., “Effect of Secondary Phases on the Properties of B4C-A1 Composites.”, Anon. Abst. of the 38th Annu. Pacific Coast Regional Meeting American Ceramic Society, American Ceramic Society, Columbus, OH, (1985) (Mechan., Prop., Experimental, 0) Chernyshova, T.A., Rebrov, A.V., “Interaction Kinetics of Boron Carbide and Silicon Carbide with Liquid Aluminium”, J. Less-Common Met., 117, 203–207 (1986) (Kinetics, Experimental, 4) Dub, S.N., Prikhna, T.A., Il’nitskaya, O.N., “Mechanical Properties of the Al-B-C Compounds Crystals” (in Russian), Sverkhtverd. Mater., 6, 12–18 (1986) (Mechan. Prop., Experimental, 22) Halverson, D.C., Pyzik, A.J., Aksay, I.A., “Boron-Carbide-Aluminum and Boron-Carbide-Reactive Metal Cermets”, US patent document 4,605,440/A/, (1986) Kisly, P.S., Prikhna T.A., Golubyak, L.S., “Properties of High-Temperature Solution Grown Aluminium Borides”, J. Less-Common Met., 117, 349–353 (1986) (Experimental, 10) Peshev, P., Gyurov, G., Khristov, M., Gurin, V.N., Korsukova, M. M., Solomkin, F.Yu., Sidorin K.K., “Preparation and some Properties of Aluminium Carboboride Single Crystals”, J. Less-Common Met., 117, 341–348 (1986) (Crys. Structure, Mechan. Prop., Optical Prop., Experimental, 16) Roszler, J.J., “Process for the Manufacture of a Material Shielding Against Neutrons” (in German), DE Patent Document 2643444/C2/, (1986) Haupt, H., Werheit, H., Siejak, V., Gurin, V.N., Korsukova, M.M., “Absorption Edge and IR-active Phonons of Al3B48C2, “Boron, Borides and Related Compounds”, Proc. 9th Int. Sympos., Werheit, H. (Ed.), Univ. Duisburg, Germany, 387–389 (1987) (Experimental, 2)


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3 [1987Kis]

[1987Lev]

[1987Pyz] [1987Sar]

[1988Kov]

[1989Blu]

[1989Hal] [1989Kha]

[1990Ase]

[1990Luk]

[1990Oka]

[1990Pyz]

[1990Ram] [1991Kha1]

[1991Kha2]

[1991Kha3]

[1991Kha4]

[1991Pri]

Al–B–C Kisly, P.S., Prikhna, T.A., Gontar A.N., Podarevskaya, O.V., “Structure and Properties of Monocrystals of the Al-B-C System Compounds” in “Boron, Borides and Related Compounds”, Proceedings 9th Int. Sympos., Werheit, H. (Ed), Univ. Duisburg, Germany, 273–274 (1987) (Thermodyn., Crys. Structure, Phys. Prop., Experimental, 1) Levinskas, D., “Evaluation of Boron Carbide Coatings”, Western Region American Nuclear Society Student Conference: Nuclear Technology for the Year 2000, American Nuclear Society, La Grange Park, IL., NM(USA), 68–71 (1987) (Experimental, 0) Pyzik, A.J., Aksay, I.A., “Multipurpose Boron Carbide-Aluminum Composite and its Manufacture via the Control of the Microstructure”, US patent document 4,702,7707 A/, 27, (1987) Sarikaya, M., Laoui, T., Milius D.L., Aksay, I.A., “Identification of a New Phase in the Al-B-C Ternary by High-Resolution Transmission Electron Microscopy”. Proc. 45th Ann. Meeting of the Electron Microscopy Society of America, Bailey, G.N., (Ed.), San Franc. Press, USA, 168–169 (1987) (Crys. Structure, Experimental, 4) Koval’chenko, M.S., Paustovskij, A.V., Bolejko, B.M. Zhidkov, A.V., “Laser Surface Hardening of Cermets on the Base of Boron Carbide”(in Russian), Poroshk. Metall., 5, 77–80 (1988) (Mechan. Prop., Experimental, 6) Blumenthal, W.R., Gray, G.T., “Structure-Property Characterization of Shock-Loaded B4C-Al”, Inst. Phys. Conf. Ser. No 102: Session 7, Paper Presented at Int. Conf. Mech. Prop. Materials at High Rates of Strain, Oxford, 363–370 (1989) (Experimental, 8) Halverson, D.C., Pyzik, A.J., Aksay, I.A., Snowden, W.E., “Processing of Boron Carbide-Aluminium Composites”, J. Am. Ceram. Soc., 72(5), 775–80 (1989) (Experimental, 33) Kharlamov, A.I., Duda, T.I., Lojchenko, S.V., Fomenko, V.V., “Preparation and Properties of Aluminium Boridocarbide Powder of Al8B4C7 Composition”, 12thUkrainian Republic Conference on Inorganic Chemistry, Vol. 1, Simferopol’, Ukr. SSR, 44pp. (1989) (Mechan. Prop., Experimental, 0) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. NATO Advanced Research Workshop, Manchester, U.K., 1989, published as ASI-Series, Ser. E: Appl. Sci., Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental,14) Lukas, H.L., “Aluminium-Boron-Carbon” in “Ternary Alloys. A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams”, Petzow, G., Effenberg, G., (Eds.), Vol. 3, VCH, Weinheim, 140–146 (1990) (Review, Phase Diagram, 14) Okada, S., Kudou, K., Hiyoshi, H., Higashi, I., Hamano, K., Lundstro¨m, T., “Preparation of AlC4B24 and Al3C2B48 Crystals”, J. Int. Ceram. Soc. Jpn., 98, 1342–1347 (1991), translated from Nippon Seramikkusu Kyokai Gakujutsu Ronbunshi, 98(12), 1330–1336 (1990) (Experimental, Crys. Structure, 24) Pyzik, A.J., Williams P.D., McCombs, A., “New Low Temperature Processing for Boron Carbide/ Aluminium Based Composite Armor”, Final Report, US-Army Research Office, DAAL 0388 C0030, 1990 (Experimental, 14) Ramesh, K. T., Ravichandran, G., “Dynamic Behavior of a Boron Carbide-Aluminum Cermet: Experiments and Observations”, Mech. Mater., 10(1-2) 19–29 (1990) (Experimental, 22) Kharlamov, A.I., Loichenko, S.V., “Electronic Transport Properties of Hot-pressed Boron-rich Compounds of the Al-B-C System” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. et al (Eds.), Albuquerque, USA, 1990, AIP, New York, 94–97 (1991) (Experimental, 5) Kharlamov, A.I., Loichenko, S.V., “Investigation: The Process of Densification of Boron-Rich Compounds of the Al-B-C System” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. et al (Eds.), Albuquerque, USA, 1990, AIP, New York, 473–481 (1991) (Experimental, 2) Kharlamov, A.I., Murzin, L.M., Loichenko, S.V., Duda, T.I., “Electrical Conductivity and Seebeck Coefficient of Hot-Pressed Specimens of Aluminium Borides and Carboborides”, Sov. Powder Metall. Met. Ceram., 9(345), 770–773 (1991), translated from Poroshk. Metall., 9(345), 62–65 (1991) (Experimental, Electr. Prop., 7) Kharlamov, A.I., Duda, T.I., Fomenko, V.V., “Preparation and Properties of High-Dispersive Powders of Aluminium Dodecaboride and Carboborides” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. (Eds.), Albuquerque, USA, 1990, AIP, New York, 512–515 (1991) (Experimental, 0) Prikhina, T.A., Kisly, P.S., “Aluminium Borides and Carboborides” in “Boron-Rich Solids”, AIP Conf. Proc. 231, Emin, D. et al (Eds.), Albuquerque, USA, 1990, AIP, New York, 590–593 (1991) (Experimental, 11)

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Al–B–C [1991Pyz] [1992Bei] [1992Var] [1992Via] [1993Bau]

[1993Gon]

[1993Ips] [1993Kau] [1993Wen] [1993Wer]

[1994Dus] [1994Kud]

[1995Bon] [1995Hil] [1995Osc] [1995Pyz]

[1995Rug] [1996Bid] [1996Hil1]

[1996Hil2]

[1996Kas] [1996Pyz] [1996RMi]

[1996Vin]

3

Pyzik, A.J., Nilson, R.T., “B4C/A1 Cermets and Method for Making Same”, US Patent Document 5,039,633, (1991) Beidler, C.J., Hauth III, W.E., Goel, A., “Development of a B4C/A1 Cermet for Use as an Improved Structural Neutron Absorber”, J. Testing and Evaluation, 20(1), 67–70 (1992) (Experimental, 6) Vardiman, R.G., “Microstructures in Aluminium, Ion Implanted with Boron and Heat Treated”, Acta Metall. Mater., 40, 1029–35 (1992) (Crys. Structure, Eperimental, 7) Viala, J.C., Gonzales, G., Bouix, J., “Composition and Lattice Parameters of a New Aluminium-Rich Boron Carbide”, J. Mater. Sci. Lett., 11, 711–714 (1992) (Crys. Structure, Experimental, 9) Bauer, J., Bittermann, H., Rogl, P., “Phase Relations and Structural Chemistry in the Ternary System Aluminium - Boron - Carbon”, COST-507, Annual Report, (1993) (Crys. Structure, Phase Diagram, Experimental, 12) Gonzalez, G., Esnouf, C., Viala, J.C., “Structural Study of a New Aluminium Rich Borocarbide Formed by Reaction at the B4C/Al Interface”, Mater. Sci. Forum, 126-128, 125–128 (1993) (Crys. Structure, Experimental, 4) Ipser, H., privat communication (1993) (Experimental) Kaufmann, L., private communication (1993) (Thermodyn.) Wen, H., “Thermodynamic Calculations and Constitution of the Al-B-C-N-Si-Ti System” (in German), Thesis, Univ. Stuttgart, 1–183 (1993) (Calculation, Phase Diagram, Thermodyn., 223) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi (B), B179, 489–511 (1993) (Crys. Structure, Experimental, 51) Duschanek, H., Rogl, P., “The System Al-B”, J. Phase Equilib., 15(5), 543–52 (1994) (Crys. Structure, Phase Diagram, Experimental, #, 78) see also ibid, 16(1), 6 (1995) Kudou, K., Okada, S., Hikichi, H., Lundstro¨m, T., “Preparation and Properties of Si-doped Al3C2B48Type Crystals” (in Japanese), J. Soc. Mater. Sci., Jpn., 43(485), 223–228 (1994) (Experimental, Crys. Structure, Phys. Prop., 20) Bond, G.M., Inal, O.T., “Shock-Compacted Aluminium/Boron Carbide Composites”, Compos. Eng. 5(1), 9–16 (1995) (Experimental, 18) Hillebrecht, H., Meyer, F., “B48A13C2 - a Filled Variant of Tetragonal Boron”, Z. Kristallogr., Suppl. 10, 101 (1995) (Crys. Structure, Experimental, 2) Oscroft, R.J., Roebuck P.H.A., Thompson, D.P., “Characterisation and Range of Composition for Al8B4C7”, Br. Ceram. Trans., 94(1), 25–26 (1995) (Experimental, 11) Pyzik, A.J., Beaman, D.R., “Al-B-C Phase Development and Effects on Mechanical Properties of B4C/ Al-Derived Composites”, J. Am. Ceram. Soc., 78(2), 305–312 (1995) (Crys. Structure, Mechan. Prop., Experimental, 25) Ruginets, R., Fischer, R. “Microwave Sintering of Boron Carbide Composites”, Am. Ceram. Soc. Bull., 74(1), 56–58 (1995) (Experimental) Bidaud, E., research at Univ. Wien, unpublished (1996) Hillebrecht, H., Meyer, F.D., “The Structure of B48Al3C2 - A Filled and Distorted Variant of Tetragonal Boron (I)” in “Boron, Borides and Related Compounds”, Proc. 12th Int. Symp., Baden/Wien, paper PA.4, 59 (1996) (Crys. Structure, Experimental, 6) Hillebrecht, H., Meyer, FD., “Synthesis, Crystal Structure, and Vibrational Spectra of Al3BC3, a Carbidecarboborate of Aluminium with Linear (C=B=C)5– Anions”, Angew. Chem., 35(21), 2499–2500 (1996), translated from Angew. Chemie, 108(21), 2655–2657 (1996) (Crys. Structure, Experimental, 17) Kasper, B., “Phase Equilibria in the B-C-N-Si System”, Thesis, Max Plank Institute-PML, Stuttgart, (1996) (Phase Diagram, Thermodyn.) Pyzik, A.J., Deshmukh, U.V., Dunmead, S.D., Ott, J.J., Allen, T.L., Rossow, H.E., “Light Weight Boron Carbide/Aluminium Cerments”, United States Patent: 5,521,016, (1996) R’Mili, M., Massardier, V., Merle, P., Vincent, H., Vincent, C., “Mechanical Properties of T300/A1 Composites. Embrittlement Effects due to a B4C Coating”, J. Mater. Sci., 31, 4533–4539 (1996) (Mechan. Prop., Experimental, 12) Vincent, H., Vincent, C., Berthet, M. P., Bouix, J., Gonzalez, G., “Boron Carbide Formation from BCl3CH4-H2 Mixtures on Carbon Substrates and in a Carbon-Fibre Reinforced Al Composite”, Carbon, 34 (9), 1041–1055 (1996) (Crys. Structure, Mechan. Prop., Experimental, 25)


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3 [1997Du] [1997Mey] [1997Sch]

[1997Via] [1997Yue] [1998Kas]

[1998Kha]

[1998Rog]

[1999Bur] [1999Cha]

[1999Tor]

[1999Tsu] [2000Hal]

[2000Hig] [2000Kha]

[2000Liu] [2000Mey]

[2000Pyz] [2000Sav] [2000Sol] [2000Wan]

Al–B–C Du, W.F., Watanabe, T., “High-Toughness B4C-AlB12 Composites Prepared by Al Infiltration”, J. Eur. Ceram. Soc., 17, 879–884 (1997) (Mechan. Prop., Experimental, 15) Meyer, F.D. Hillebrecht, H., “Synthesis and Crystal Structure of Al3BC, the First Boridecarbide of Aluminium”, J. Alloy. Compd., 252, 98–102 (1997) (Crys. Structure, Experimental, 30) Schmechel, R., Werheit, H., Robberding, K., Lundstro¨m, T., Bolmgren, H., “IR-active Phonon Spectra of B-C-Al Compounds with Boron Carbide Structure”, J. Solid State Chem., 133, 254–259 (1997) (Experimental, 11) Viala, J.C., Bouix, J., Gonzalez, G., Esnouf, C. “Chemical Reactivity of Aluminium with Boron Carbide”, J. Mater. Sci, 32, 4559–4573 (1997) (Phase Diagram, Experimental, 39) Yu¨cel, O., Tekin, A., “The Fabrication of Boron-Carbide-Aluminium Composites by Explosive Consolidation”, Ceram. Int., 23, 149–152 (1997) (Experimental, Mechan. Prop., 3) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Kharlamov, A.I., Kirillova, N.V., Loichenko, S.V., Fomenko, V.V., “Properties of Aluminium Borides and Borocarbides”, Russ. J. Appl. Chem., 71(5), 743–749 (1998), translated from Zh. Prikl. Khim, 71(5), 717–724 (1998) (in Russian), (Crys. Structure, Kinetics, Mechan. Prop., Experimental, 13) Rogl, P., “Al-B-C (Aluminium-Boron-Carbon)”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12170.2.20, (1998) aslo published in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G., (Ed.), ASM-Intl, MSI, 3–15 (1998) (Assessment, Crys. Structure, Experimental, Phase Diagram, 50) Burkhardt, U., Grin, Y., “Refinement of the Aluminium Diboride Crystal Structure” in “Borides and Related Compounds”, Abst. 13th Int. Symp. on Boron, Dinar (France), 13pp., (1999) (Crys. Structure, 3) Chapman, T.R., Niesz, D.E., Fox, R.T., Fawcett, T., “Wear-resistant Aluminum - Boron - Carbide Cermets for Automotive Brake Applications”, Wear, 236, 81–87 (1999) (Mechan. Prop., Experimental, 9) Torquato, S., Yeong, C.L.Y., Rintoul, M.D., Milius, D.L., Aksay, I.A., “Elastic Properties and Structure of Interpenetrating Boron Carbide/Aluminum Multiphase Composites”, J. Am. Ceram. Soc., 82(5), 1263–1268 (1999) (Mechan. Prop., 32) Tsuchida, T., Kan, T., “Synthesis of Al3BC in Air from Mechanically Activated Al/B/C Powder Mixtures”, J. Eur. Ceram. Soc., 19, 1795–1799 (1999) (Crys. Structure, Experimental, 12) Hall, A., Economy, J., “The Al(L)+AlB12ÐAlB2 Peritectic Transformation and its Role in the Formation of High Aspect Ratio AlB2 Flakes”, J. Phase Equilib., 21(1), 63–69 (2000) (Phase Diagram, Experimental, 21) Higashi, I., “Crystal Chemistry of α-AlB12 and γ-AlB12”, J. Solid State Chem., 154, 168–176 (2000) (Crys. Structure, Experimental, 18) Kharlamov, A. I., Nizhenko, V.I., Kirillova, N.V., Floka, L.I., “Wettability of Hot-Pressed Samples of Boron-Containing Aluminium Compounds by Liquid Metals and Alloys” (in Russian), Zh. Prikl. Khim., 73(6), 884–888 (2000) (Experimental, 14) Liu, C.H., “Structure and Properties of Boron Carbide with Aluminum Incorporation”, Mater. Sci. Eng. B, B72, 23–26 (2000) (Phys. Prop., Crys. Structure, Experimental, 10) Meyer, F.D., Hillebrecht, H., “Ternary Phases in the System Al/B/C” in “High Temperature Materials Chemistry”, Vol. 15, Part 1, K. Hilpert et al. (Eds.), Proc. 10th Intl. IUPAC Conf., Forschungszentrum Ju¨lich, Germany, Published by Schriften des Forschungszentrums Juelich, 161–164 (2000) (Crys. Structure, 5) Pyzik, A.J., Deshmukh, U.V., Krystosek, R. D., “Aluminum-Boron-Carbon Abrasive Article and Method to Form Said Article”, US Patent: 6,042,627, (2000). Savyak, M., Uvarova, I., Timofeeva, I., Isayeva L., Kirilenko, S., “Mechanochemical Synthesis in Ti-C, Ti-B, B-C, B-C-A1 Systems”, Mater. Sci. Forum, 343-346, 411–416 (2000) (Experimental, 4) Solozhenko, V.L., Meyer, F.D., Hillebrecht, H., “300-K Equation of State and High-Pressure Phase Stability of Al3BC3”, J. Solid State Chem., 154, 254–256 (2000) (Crys. Structure, Experimental, 11) Wang, T., Yamaguchi, A., “Some Properties of Sintered Al8B4C7”, J. Mater. Sci. Letter., 19, 1045–1046 (2000) (Calculation, Crys. Structure, 6)

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Al–B–C [2000Wer] [2001Fje]

[2001Lee]

[2002Ars] [2002Bur]

[2002Kou1] [2002Kou2]

[2002Zhe] [2003Per]

[Mas2] [V-C2]

3

Werheit, H., Schmechel, R., Meyer, F.D., Hillebrecht, H., “Interband Transitions and Optical Phonons of B48Al3C2”, J. Solid State Chem., 154, 75–78 (2000) (Optical Prop., Experimental, 10) Fjellstedt, J., Jarfors, A.E.W., El-Benawy, T., “Experimental Investigation and Thermodynamic Assessment of the Al-rich Side of the Al-B System”, Mater. Des., 22(6), 443–449 (2001) (Thermodyn, Phase Diagram, Experimental, 14) Lee, K.B., Sim, H.S., Cho, S.Y., Kwon, H., “Reaction Products of Al-Mg/B4C Composite Fabricated by Pressureless Infiltration Technique”, Mater. Sci. Eng. A, 302, 227–234 (2001) (Crys. Structure, Phase Diagram, Experimental, 17) Arslan, G., Kara, F., Turan, S., “Mechanical Properties of Melt Infiltrated Boron Carbide-Aluminium Composites”, Key Eng. Mater., 206-213(2), 1157–1160 (2002) (Experimental, Mechan. Prop., 5) Burkhardt, U., Gurin, V., Borrmann, H., Schnelle, W., Grin, Y., “On the Electronic and Structural Properties of Aluminium Diboride Al0.9B2” in “Boron, Borides and Related Compound”, Abst. 14th Int. Symp., (ISBB’02), Saint Petersburg, O4, (2002) (Crys. Structure, 3) Kouzeli, M., Mortensen, A., “Size Dependent Strengthening in Particle Reinforced Aluminium”, Acta Mater., 50, 39–51 (2002) (Mechan. Prop., Experimental, 59) Kouzeli, M., Marchi, C. S., Mortensen, A., “Effect of Reaction on the Tensile Behavior of Infiltrated Boron Carbide-Aluminum Composites”, Mater. Sci. Eng. A, A337, 264–273 (2002) (Experimental, Mechan. Prop., 51) Zheltov, P., Grytsiv, A., Rogl, P., Velikanova, T.Ya., Research at Univ. Wien (unpublished) (2002) (Phase Diagram, Crys. Structure) Perrot, P., “Aluminium-Carbon”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, to be published, (2003) (Phase Diagram, Crys. Structure, Assessment, 19) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)


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Aluminium – Boron – Molybdenum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Qingsheng Ran, updated by Lazar Rokhlin, Tatiana Dobatkina, Elena Semenova, Natalia Kol’chugina

Introduction The Al-B-Mo system is of interest mainly because of the application of low-alloyed Mo compositions in manufacturing parts for operation under conditions of high-temperature gases and erosion wear [1988Buk, 1994Buk, 1995Buk, 1997Buk, 1999Buk, 2003Buk1, 2003Buk2]. The first evaluation within the ongoing MSIT Evaluation Program was made by [1990Ran], which is updated by the present work. An isothermal section at 1000˚C was presented by [1965Rie] obviously from the determination of phase regions using X-ray diffraction. Metals were cold pressed; samples were prepared by reacting at 900˚C and annealing at 1200˚C for 15 h. While cooling, equilibrium was said to be frozen in at 1000˚C. [1942Hal] prepared a ternary compound “Mo7Al6B7” by thermal reaction of MoO3, B2O3, Al and S and attempted a first indexation of its powder pattern. The existence of the phase was later confirmed by [1965Rie, 1966Jei], which assumed that its formula is MoAlB. Moreover, an Al-stabilized molybdenum boride with the CrB type structure was found. To investigate the structure of this phase more accurately, [1966Jei] heated a sample of the composition Mo:Al:B = 1:6:1 to about 1800˚C and cooled it for 40 minutes to 1000˚C. The charge was subsequently treated in hot NaOH solution and single crystals were obtained from the residue. In performing X-ray diffraction, the Debye-Scherrer, Weissenberg, and rotating crystal methods were used. In [1977Gur], crystals of solid solutions of boron (to 1 mass%) in MoAl4(5) were separated from Al-B-Mo alloys (the alloys were prepared preliminarily by melting followed by spontaneous solidification) by dissolution of Al matrix in HCl. The MoAlB (orthorhombic) compound and another compound of the Al-B-Mo system with a rhombohedral crystal structure were reported by [1976Hig]. Reviews [1972Cha, 1976Hig] included the aforementioned data on the phases existing in the ternary Al-B-Mo system. Investigations into the Al-B-Mo phase relations, structure identifications are given in Table 1.

Binary Systems The binary systems Al-B [2004Mir] (Fig. 1), B-Mo [1992Rog] (Fig. 2), and Al-Mo [2005Sch] are accepted in the present assessment. Landolt‐Bo¨rnstein New Series IV/11E1

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Solid Phases The solid phases of the Al-B-Mo system are described in Table 2 that is presented mainly with allowance for the binary phase diagrams accepted in this assessment. Two ternary phases were found in the system; these are τ1 (MoAlB) and τ2 (Mo9AlB10). The homogeneity ranges of the phases are taken from the isothermal section at 1000˚C given by [1965Rie]. The aluminium solubility in B-Mo phases and boron solubility in Al-Mo phases are taken also from this section. Virtually the same boron solubility in MoAl4 was confirmed by [1977Gur].

Isothermal Sections The only isothermal section at 1000˚C (Fig. 3) is constructed based on data by [1965Rie] corrected with allowance for the binary phase diagrams accepted in the present evaluation. In accordance with the binary B-Mo system, the existing MoB4 compound is added in Fig. 3; two three-phase regions containing the MoB4 phase in the section are given tentatively. The MoAl4 and Mo4Al17 compounds are added to the section after the accepted binary Al-Mo system. They are located between the Al rich liquid and Mo3Al8 that are in equilibrium with the τ1 phase. Therefore, it was reasonable to accept the existence of the MoAl4 and Mo4Al17 compounds in equilibrium with the τ1 phase. This is indicated in the section in Fig. 3.

Notes on Materials Properties and Applications Owing to the importance of molybdenum-based alloys containing small amounts of Al and B for the manufacturing of parts for operation under conditions of high-temperature gases and erosion wear, various mechanical properties of them have been investigated at room and high temperatures [1988Buk, 1994Buk, 1995Buk, 1997Buk, 1999Buk, 2000Buk, 2003Buk1, 2003Buk2]. Table 3 shows the mechanical properties and their correlation with the structural state.

Miscellaneous The possibility of forming composites with Al matrix strengthened by molybdenum borides is considered in [1996Bie].

. Table 1 Investigations of the Al-B-Mo Phase Relations, Structures and Thermodynamics Reference [1942Hal]

Method/Experimental Technique

Temperature/Composition/Phase Range Studied

Interaction of molybdenum boride with 74 mass% Mo, 18 mass% Al, 8.5 aluminium followed by dissolution of Al mass% B / near Mo7Al6B7 matrix / X-ray diffraction, chemical analysis, density measurements

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. Table 1 (continued) Reference


Temperature/Composition/Phase Range Studied 1000˚C / 0–80 at.% Al, 0–100 at.% Mo, 0–100 at.% B /

[1965Rie]

Cold pressing of powders followed by heating, X-ray diffraction

[1966Jei]

Heating of components to 1800˚C followed 1000˚C / MoAlB by slow cooling, chemical extraction / X-ray diffraction

[1977Gur] Melting followed by chemical extraction after solidification / X-ray diffraction, chemical analysis

1400–1000˚ / 55 mass% Al, 43 mass% Mo, 0.7 mass% B / MoAl4-based solid solution of B





Comments/References

(βAl)

hP2 P63/mmc Mg

a = 269.3 c = 439.8

at 25˚C, 20.5 GPa [Mas2]

(αAl) < 660.452

cF4 Fm 3m Cu

a = 404.96

at 25˚C [Mas2]

(βB) < 2092

hR333 R3m βB

a = 1093.30 c = 2382.52

[1993Wer]

a = 1093.02 c = 2381.66

pure B [1976Lun]

a = 1006.1 c = 1421.0

presumably metastable phase, possibly stabilized by impurities, preparation at 1100 to 1400˚C [1971Amb]

(B) (tetr.)

tP192 derived from αAlB12

(αB)

hR36 R 3m αB

(Mo) < 2623


cI2 Im 3m W

a = 490.8 c = 1256.7 a = 314.70

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presumably metastable phase, preparation below 1000˚C [1971Amb] pure B, single crystal [1994Cha] at 25˚C [Mas2] dissolves up to 19.5 at.% Al and up to 1 at.% B [2005Sch, Mas2]

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Comments/References

αAlB12 < 2094

oP384 P212121 αAlB12

a = 1662.3 b = 1754.0 c = 1018.0

[2004Mir] [V-C2]

γAlB12 ≤ 1550

oP396 P212121 γAlB12

a = 1662.3 b = 1754.0 c = 1018.0

[1983Hig] Metastable or metal impurity stabilized [1994Dus]

AlB2 < 972 - 213

hP3 P6/mmm AlB2

a = 300.5 c = 325.7

[2004Mir] [V-C2]

MoAl12 < 712

cI26 Im 3 WAl12

a = 757.3

[2005Sch] 92.3 at.% Al

MoAl5(h2) 846 - (750 - 800)

hP12 P6322 WAl5

a = 491.2

[2005Sch] 83.3 at.% Al

a = 493.3 c = 4398

[2005Sch] 83.3 at.% Al

MoAl5(h1) hP60 (750–800) - 648 P321 MoAl5(h1) MoAl5(r) ≲ 648

hP36 R 3c MoAl5(r)

a = 495.1 c = 2623

[2005Sch] 83.3 at.% Al

Mo5Al22(h) 964 - 831

oF216 Fdd2 Mo5Al22(h)

a = 7382 ± 3

[2005Sch] 81.5 at.% Al

Mo4Al17 < 1034

mC84 C2 Mo4Al17

a = 915.8 ± 0.1 b = 493.23 ± 0.08 c = 2893.5 ± 0.5 β = 96.71 ± 0.01˚

[2005Sch] 81 at.% Al

MoAl4(h) 1177 - 942

mC30 Cm WAl4

a = 525.5 ± 0.5 [2005Sch] b = 1776.8 ± 0.5 79 at.% Al c = 522.5 ± 0.5 β = 100.88 ± 0.06˚

Mo1–xAl3+x 1260 - 1154

cP8 Pm 3n Cr3Si

a = 494.5 ± 0.1

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[2005Sch] 76.5 at.% Al


Al–B–Mo

4





Comments/References

MoAl3(h) 1222 - 818

mC32 C2/m MoAl3

a = 1639 ± 1.0 [2005Sch] b = 359.4 ± 0.1 75 at.% Al c = 838.6 ± 0.4 β = 101.88 ± 0.07˚

Mo3Al8 < 1555 ± 10

mC22 Cm Mo3Al8

a = 920.8 ± 0.3 [2005Sch] b = 363.78 ± 0.03 72.7 at.% Al c = 1006.5 ± 0.3 β = 100.78 ± 0.05˚

ζ1(h), Mo2Al3(h) 1570 - 1490

Unknown

-

[2005Sch] 63 at.% Al

ζ2(h), MoAl(h) 1750 - 1535

cP2 Pm 3m CsCl

a = 309.8

[2005Sch] 50–52 at.% Al

Mo3Al ≲ 2150

cP8 Pm 3n Cr3Si

a = 495.0

[2005Sch] 23–29 at.% Al

Mo2B < 2280

tI12 I4/mcm Al2Cu

a = 554.7 c = 473.9

[Mas2] [V-C2] 33 at.% B

αMoB < 2180

tI16 I41/amd MoB

a = 310.3 c = 1695

[Mas2] [V-C2] 48–50 at.% B

βMoB 2600 - 2180

oC8 Cmcm CrB

MoB2 2375 - 1517

hP3 P6/mmm AlB2

a = 303.9 c = 305.5

[Mas2] [V-C2] 62–66 at.% B

Mo2B5 < 2140

hR21 R3m Mo2B5

a = 300.7 c = 2091.0

[Mas2] [V-C2] 67–69 at.% B

MoB4 < 1807

hP20 P63/mmc WB4

a = 520.3 c = 634.5

[Mas2] [V-C2] 79 at.% B


a = 315.1 b = 847.0 c = 308.2

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[Mas2] [V-C2] 48–51 at.% B

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Comments/References

* τ1, MoAlB ≲ 1000˚C

oC12 Cmcm BCU

a = 321.2 b = 1398.5 c = 310.2

[1966Jei, V-C2] 31–34.5 at.% B and 33–35.5 at.% Al

* τ2, Mo9AlB10

oC8 Cmcm CrB

a = 316.3 b = 847.5 c = 308.2

[V-C2] 48–50 at.% B and 5–6.5 at.% Al

. Table 3 Investigations of the Al-B-Mo Materials Properties Reference

Method / Experimental Technique

Type of Property

[1988Buk]

Tensile creep tests at 1500–2000˚C. Creep rate and stress-rupture strength for Mobased compositions low-alloyed with Al and B

[1994Buk]

Tensile creep tests 1500–2000˚C.

Correlation between short-term, long-term, and low-cycle strength characteristics of Mo-based compositions low-alloyed with Al and B were established.

[1995Buk]

Tensile tests 20–2000˚C.

Ultimate strength and relative elongation of Mo-based compositions low-alloyed with Al and B are determined.

[1997Buk]

Fatigue tests under conditions of rigid alternating bending at 950–1550˚C.

Effect of heat treatment and welding on fatigue resistance of Mo-based compositions lowalloyed with Al and B is studied.

[1999Buk]

Tensile tests at (0.5–0.8) Tm.

Correlation between short-term, long-term static, creep resistance, and few-cycle strength characteristics of Mo-based compositions lowalloyed with Al and B were established.

[2000Buk]

Tensile and fatigue tests under conditions of rigid alternating bending.

Empirical correlations between fatigue strength, conventional yield strength, and grain size were established.

[2003Buk1] Tensile tests at 20–2000˚C and fatigue tests.

Correlations between the structural state and short-term and low- and high-cycle fatigue were established.

[2003Buk2] Tensile tests at 20–2000˚C.

Effect of heat treatment on ultimate strength, yield strength, relative elongation, and relative uniform deformation is considered.

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. Fig. 1 Al-B-Mo. The Al-B binary system


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. Fig. 2 Al-B-Mo. The B-Mo binary system

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. Fig. 3 Al-B-Mo. Isothermal section at 1000˚C


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References [1942Hal] [1965Rie]

[1966Jei] [1971Amb] [1972Cha]

[1976Hig] [1976Lun] [1977Gur]

[1983Hig] [1988Buk]

[1990Ran]

[1992Rog]

[1993Wer]

[1994Buk]

[1994Cha]

[1994Dus] [1995Buk]

[1996Bie]

Halla, F., Thury, W., “On Borides of Molybdenum and Tungsten” (in German), Z. Anorg. Allg. Chem., 249, 229–237 (1942) (Crystal Structure, Experimental, 7) Rieger, W., Nowotny, H., Benesovksy, F., “Complex Borides of the Transition Metals (Mo, W, Fe, Co, Ni)” (in German), Monatsh. Chem., 96(3), 844–851 (1965) (Crys. Structure, Experimental, Phase Diagram, #, 10) Jeitschko, W., “The Crystal Structure of MoAlB” (in German), Monatsh. Chem., 97, 1472–1476 (1966) (Crys. Structure, Experimental, 10) Amberger, E., Ploog, G., “Formation of Lattices of Pure Boron” (in German), J. Less-Common Met., 23, 21–31, (1971) (Crys. Structure, Experimental, 17) Chaban, N.F., Kuz’ma, Yu.B., “Metallides in the Period IV Transition Metal-Aluminum (Gallium) Boron Ternary System” (in Russian), Stroenie, Svoistva i Primenenie Metall., 102–107 (1972) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, 18) Higashi, I., Takahashi, Y., Atoda, T., “Crystal Growth of Borides and Carbides of Transition Metals from Al Solutions”, J. Cryst. Growth, 33, 207–211 (1976) (Experimental, 16) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Gurin, V.N., Korsukova, M.M., Popov, V.I., Elizarova, O.V., Belozov, N.N., Kuz’ma, Yu.B., “Solid Solutions of B, C, Si in Aliminides of Transient Metals” (in Russian), Akad. Nauk SSSR, Nauka, Moscow, 39–42 (1977) (Crys. Structure, Experimental, Phase Diagram, 6) Higashi, I., “Aluminum Distribution in the Boron Framework of γ-AlB12”, J. Solid State Chem., 47, 333–349 (1983) (Experimental, Crys. Structure, 21) Bukhanovskii, V.V., Kharchenko, V.K., Polishchuk, E.P., Kravchenko, V.S., Galinzovskaya, T.D., Onoprienko, A.A., “Influence of Production Factors on the High-Temperature Strength Characteristics of Molybdenum Alloys”, Strength Met., 20(6), 818–824 (1988), translated from Probl. Prochn., (6), 102–108 (1988) (Experimental, Kinetics, Morphology, 18) Ran, Q., “Al-B-Mo (Aluminium-Boron-Molybdenum)” in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.19496.1.20, (1990), also published in “Ternary Alloys”, Petzow, G., Effenberg, G. (Eds.), Vol. 3, VCH Verlagsgesellschaft, Weinheim, Germany, 189–191 (1990) (Phase Diagram, Phase Relations, Review, 5) Rogl, P., “The system B-N-Mo” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), Ohio, USA: ASM, Materials Park, 64–67 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Experimental, Review, *, 10) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Stat. Sol. B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) Bukhanovskii, V.V., Borisenko, V.A., Kharchenko, V.K., “Correlations Between Short-Term, Long-Term Static, and Low-Cycle Strength Characteristics of Low-Alloy Molybdenum Alloys at High Temperatures”, Strength Met., 28(12), 895–902 (1994), translated from Probl. Prochn., (12), 43–51 (1994) (Experimental, Mechan. Prop., 21) Chakrabarti, D.J., Laughlin, D.E., “B-Cu (Boron-Copper)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 74–78 (1994) (Review, Phase Diagram, Crys. Structure, Thermodyn., 24) Duschanek, H., Rogl, P., “The Al-B (Aluminum-Boron) System”, J. Phase Equilib., 15, 543–552 (1994) (Assessment, Crys. Structure, Phase Relations, Phase Diagram, Review, Thermodyn., 78) Bukhanovskii, V.V., Borisenko, V.A., Kharchenko, V.K., “Mechanical Characteristics of Molybdenum Alloys of the Systems the Mo-Al-B and Mo-Zr-B in the Temperature Range 290–2270 K”, Strength Met., 27(11-12), 688–695 (1995), translated from Probl. Prochn., 27(11-12), 70–80 (1995) (Experimental, Mechan. Prop., 19) de Bie, J.E., Froyen, L., Lust, P., Delaey, L., “Solidification of In-Situ Aluminium Composites”, Mater. Sci. Forum, 215-216, 435–442 (1996) (Morphology, Phase Relations, Theory, 10)

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Al–B–Mo [1997Buk]

[1999Buk] [2000Buk]

[2003Buk1]

[2003Buk2]

[2004Mir]

[2005Sch]

[Mas2] [V-C2]

4

Bukhanovskii, V.V., Borisenko, V.A., Kharchenko, V.K., “Effect of Heat Treatment and Welding on the Fatigue Resistances of Molybdenum Alloys of the Mo-Zr-B and Mo-Al-B Systems”, Met. Sci. Heat Treat., 39(5-6), 247–250 (1997), translated from Metalloved. Term.. Obrab. Met., (6), 16–19 (1997) (Experimental, Mechan. Prop., 8) Bukhanovskii, V.V., “Correlation Between the Strength and Creep of Molybdenum and Tungsten Alloys”, Russ. Metall. (Engl. Transl.), (5), 91–98 (1999) (Experimental, Mechan. Prop., 22) Bukhanovskii, V.V., “Correlations between the Characteristics of Fatigue Resistance, Short-Term Strength, and Structure of Low-Molybdenum Alloys”, Strength Met., 32(4), 361–367 (2000), translated from Probl. Prochn., (4), 75–85 (2000) (Experimental, Mechan. Prop., 14) Bukhanovsky, V., Mamuzic, I., Borisenko, V., “Interrelation between the Structural State of Material and Mechanical Properties of Low-Alloyed Molybdenum Alloys”, Metallurgia, 42(3), 159–166 (2003) (Experimental, Kinetics, Mechan. Prop., 25) Bukhanovsky, V.V., Mamuzic, I., Borisenko, V.A., “The Effect of Thermal Treatment on the Mechanical Properties of Low-Alloyed Molybdenum Alloys Over the Wide Range of Temperatures”, Metallurgia, 42(1), 9–14 (2003) (Experimental, Morphology, Mechan. Prop., 20) Mirkpvic, D., Gro¨bner, J., Schmid-Fetzer, R., Fabrichnaya, O., Lukas, H.L., “Experimental Study and Thermodynamic Re-Assessment of the Al-B System,” J. Alloys Compd., 384, 168–174 (2004) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calculation, #, 26) Schuster, J.C., “Al-Mo (Aluminum-Molybdenum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 30.12123.1.20, (2005) (Crys. Structure, Phase Diagram, Assessment, 61) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)


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Aluminium – Boron – Silicon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Hans Leo Lukas

Introduction The first evaluation within the ongoing MSIT Evaluation Program was made by [1990Luk], which is updated by the present work. The only experimental ternary information on this system is the existence of a ternary phase Al3SiB48 [1969Lam, 1972Lam] (Table 1). There is evidence that this is a high temperature phase, which is easily quenchable to room temperature by rapid or moderate cooling. The previously reported βAlB12 phase is assumed to be identical with this phase, the structure of which is closely related to that of pure B. A phase with the same structure also appears as Al3B48C2 [1964Mat]. [1981Por] estimated interaction parameters of B and Si in the (Al) solid solution. [1982Doe] calculated the whole ternary phase diagram based on assessments of the binary subsystems. Since that time newer assessments for all three binary systems are available. Especially for the Al-B system good calorimetrically measured enthalpy data are available for the intermediate phases [2001Mes] increasing the quality of the new assessment of [2004Mir]. The diagrams given here are from calculations using the dataset of [2005Gro], which is composed of the binary assessments of Al-B [2004Mir], Al-Si [1996Gro] and B-Si [1998Fri]. They differ markedly from those of [1982Doe] although the reason is only a slightly more negative Gibbs energy of mixtures of the same overall compositions of AlB12 + Si compared to those of liquid-Al + B-Si phases, whereas in the dataset of [1982Doe] the contrary was assumed. The tie-lines between AlB12 and Si-Al melt, as calculated by [2005Gro], agree with experimental findings of [2005Yos1] [1999Wan, 2000Wan, 2003Nog, 2005Hui] investigated the influence of B-additions in the magnitude of 100 to 1000 mass-ppm to the morphology of the eutectic microstructure of hypoeutectic Al-Si alloys. The two newest ones did not find significant effects of B.

Binary Systems For the thermodynamic calculations the following assessments were merged by [2005Gro] to a ternary dataset: Al-B [2004Mir], Al-Si [1996Gro] and B-Si [1998Fri]. The phase diagrams calculated from these binary datasets agree well with those compiled in [Mas2], except Al-B, where [Mas2] assumes an additional phase AlB10, which was proved to be stable only as carbon containing ternary phase [1964Mat, 1994Dus].


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Solid Phases One ternary phase, Al3SiB48, was reported by [1969Lam]. The authors got either this phase or αAlB12 on heating samples between 1400 and 1600˚C, indicating the upper temperature limit to be in this range. This phase was formerly described as binary phase βAlB12, however, [1964Mat] proved, that it is stable only with ternary additions, and gave a formula Al3B48C2. [1969Lam] showed, that carbon can be replaced in this phase by silicon, but not by germanium. The only stable binary Al-B phases are AlB2 and αAlB12 [1964Mat, 1994Dus]. Three more stable binary intermediate phases exist in the B-Si system, see Table 2. The crystal structures of the unary phases of pure boron are differently described in literature [V-C2]. The most likely interpretation is the existence of a phase hR36 below ca. 1100˚C and a high temperature phase B described by the Pearson symbols hR105, hR111 or hR141 regarding the number of atomic positions per rhombohedral unit cell. The hexagonal lattice parameters agree fairly well, a = 1092–1096, b = 2381–2389 pm. Regarding the partial occupations of the different structure descriptions, the number of atoms per rhombohedral unit cell is nearly equal for the different descriptions, 105–107. All well characterized solid phases are summarized in Table 2. Some of the crystal structures contain sites occupied only partially, there two Pearson symbols are given, one counting the sites, the other one counting the atoms per unit cell.

Invariant Equilibria Figure 1 shows the reaction scheme derived from calculations using the dataset of [2005Gro]. The calculated temperatures and phase compositions are given in Table 3. The scheme and especially the temperatures have to be taken as tentative. The temperature limits of Al3SiB48 are not known, there is only evidence that it forms between 1400 and 1600˚C and that it transforms at lower temperatures to another phase for which no details are known. Phase Al3SiB48 (τ) was added to the dataset of [2005Gro] as stoichiometric phase with linearly temperature dependent Gibbs energy forming peritectoidally at 1500˚C and neglecting the possible transformation. Below 1500˚C this phase appears everywhere in equilibrium with the three phases AlB12, SiB36 and SiB6.

Liquidus, Solidus and Solvus Surfaces The liquidus surface resulting from the calculation is shown in Fig. 2. It is remarkably different from that calculated from the dataset of [1982Doe] due to a newer assessment of the Al-B system [2004Mir] based on precise calorimetric measurements of the enthalpies of the AlB intermediate phases [2001Mes]. The isotherms are less dependent on the thermodynamic quantities than the curves of double saturation and therefore they should be taken as reasonably well defined.

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Isothermal Sections The calculated isothermal sections at 1300 and 500˚C are shown in Figs. 3 and 4 respectively.

Temperature – Composition Sections Figure 5 shows a partial calculated Temperature - Composition section at 0.1 at.% B. Note that here the field L+AlB2+AlB12 is so narrow, that it is degenerated to a single line.

Thermodynamics There are no experimental data on ternary thermodynamic quantities. Some binary data, however, are very important for ternary calculations: the calorimetrically measured enthalpies of formation of AlB2 and AlB12 [2001Mes] as well as the enthalpy change during the invariant reaction L + AlB12 Ð AlB2 [2004Mir] A thermodynamic dataset was assessed by [2005Gro] by merging the binary datasets of [2004Mir] (Al-B), [1996Gro] (Al-Si) and [1998Fri] (B-Si) without using ternary terms.

Notes on Materials Properties and Applications Wang et al. [1999Wan, 2000Wan] reported a good refining effect of B on hypoeutectic Al-Si alloys, but [2003Nog] and [2005Hui] could not find a significant effect of B on the microstructure of the eutectic in Al-Si alloys. [2005Yos2] demonstrated, that “temperature gradient zone melting” (TGZM) with a thin Al-Si-melt may help to remove B-impurities from semiconductor grade Si material.

. Table 1 Investigations of the Al-B-Si Phase Relations, Structures and Thermodynamics Reference



[1969Lam], Single-crystal X-ray diffraction [1972Lam]

Identification of the ternary phase τ, Al3SiB48

[2005Yos1] Distribution coefficient of B between Al-Si melt and solid Si

1100–1300˚C


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. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C]


(Al) < 660.452

cF4 Fm3m Cu

a = 404.96

pure Al at 25˚C [V-C2]

(βB)(h) < 2092

hR315(a,b)

a = 1096 c = 2389

hR333(a), hR315(b)

a = 1093 c = 2381 a = 1093 c = 2382

three different complete structure descriptions, cited including atomic positions by [V-C2]

hR423(a), hR321(b) R3m βB


Comments/References

(αB)(r)

hR36(c) R3m αB

a = 491.1 c = 1257.3

[V-C2] presumably metastable phase

(Si) < 1414

cF8 Fd3m C diamond

a = 543.06

at 25˚C, [V-C2]

γAlB12

oP396(a), oP377(b) P212121 γAlB12

a = 1662.3 b = 1754.0 c = 1018.0

[1983Hig] Metastable or metal impurity stabilized [1994Dus]

αAlB12 ≲ 2094

tP216(a), tP189(b) P41212 or P43212 αAlB12

a = 1016.1 c = 1428.3

[1977Hig]

AlB2 972 - ~213

hP3(a), hP2.9(b) P6/mmm AlB2

a = 300.9 c = 326.2

[1964Mat]

SiB36 < 2137

hR339(a), hR314(b)

a = 1101 [V-C2] c = 2390 a = 1099 to 1113 93.3 to 97 at.% B [1984Ole] c = 2383 to 2400

SiB6 < 1850

oP340(a), oP314.5(b) a = 1439.7 b = 1831.8 Pnnm c = 991.1 B6Si

R3m B36Si

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. Table 2 (continued) Phase/ Temperature Range [˚C]


B3Si < 1270

hR45 R3m B4C

* τ, Al3SiB48

tP68(a), tP52(b) P42/nnm B25C ?


Comments/References

a = 632 to 635 [1984Ole] c = 1269 to 1275

(c)

a = 891 c = 505

[1969Lam] metastable at room temperature? Pearson symbol (a) and prototype [V-C2]

Pearson symbol showing number of sites per unit cell (for hR hexagonal unit cells) Pearson symbol showing number of atoms per unit cell (for hR hexagonal unit cells) (c) Number of atoms in Pearson symbol is valid for hexagonal unit cell. (a)

(b)

. Table 3 Invariant Equilibria Composition (at.%) Reaction

T [˚C]

Type

Phase

L + (βB) Ð αAlB12 + SiB36

2013

U1

L

L + SiB36 Ð αAlB12 + SiB6

αAlB12 + SiB36 + SiB6 Ð τ

L + SiB6 Ð αAlB12 + (Si)


1838

≈1500

1373

U2

P1

U3

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Al

B

Si

2.6

91.2

6.2

(B)

0.0

97.9

2.1

αAlB12

7.7

92.3

0.0

SiB36

0.0

96.7

3.3

L

1.9

63.0

35.1

SiB36

0.0

94.1

5.9

αAlB12

7.7

92.3

0.0

SiB6

0.0

86.2

13.8

αAlB12

7.7

92.3

0.0

SiB36

0.0

94.1

5.9

SiB6

0.0

86.2

13.8

τ

5.8

92.3

1.9

L

2.7

8.2

89.1

SiB6

0.0

85.5

14.5

αAlB12

7.7

92.3

0.0

(Si)

0.0

1.0

99.0

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. Table 3 (continued) Composition (at.%) Reaction

T [˚C]

Type

Phase

SiB6 + (Si) Ð SiB3, αAlB12

1270

D1

SiB6

L + αAlB12 Ð AlB2 + (Si)

677

U4

B

Si

0.0

85.5

14.5

(Si)

0.0

0.7

99.3

SiB3

0.0

73.8

26.2

αAlB12

7.7

92.3

0.0

81.4

0.1

18.5

L αAlB12

L Ð (Al) + (Si), AlB2

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D2

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Al

7.7

92.3

0.0

AlB2

33.3

66.7

0.0

(Si)

0.0

0.01

99.99

L

87.92

0.03

12.05

(Al)

98.5

0.0

1.5

(Si)

0.0

0.0

100.0

AlB2

33.3

66.7

0.0


5

. Fig. 1 Al-B-Si. Reaction scheme

Al–B–Si


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. Fig. 2 Al-B-Si. Liquidus surface projection

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. Fig. 3 Al-B-Si. Isothermal section at 1300˚C


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. Fig. 4 Al-B-Si. Isothermal section at 500˚C

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. Fig. 5 Al-B-Si. Partial temperature-composition section at 0.2 at.% B


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References [1964Mat] [1969Lam]

[1972Lam]

[1977Hig] [1981Por]

[1982Doe]

[1983Hig] [1984Ole] [1990Luk]

[1994Dus] [1996Gro] [1998Fri]

[1999Wan] [2001Mes]

[2000Wan] [2003Nog] [2004Mir]

[2005Gro]

[2005Hui] [2005Yos1]

[2005Yos2]

Matkovich, V.I., Economy, J., Giese, R.F., “Presence of C in Al Borides”, J. Am. Ceram. Soc., 86, 2337–2340 (1964) (Crys. Structure, Experimental, 14) Lamikhov, L.K., Neronov, V.A., Rechkin, V.N., Samsonova, T.I., “On the β-AlB12 Phase in the Si-Al-B System” in org. Mater. (Engl. Trans.), 5, 1034–1036 (1969), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 5(7), 1214–1217 (1969) (Experimental, Morphology, 12) Lamikhov, L.K., Neronov, V.A., Rechkin, V.N., Samsonova, T.I., “β-AlB12 Phase in the Si-Al-B System” (in Russian), in “Metalloterm. Metody Poluch. Soedin. Splavov”, Kornilov, A.A. (Ed.), Nauka, Sib. Otdel. Novosibirsk, 52–57 (1972) (Review, Crys. Structure, 12) Higashi, I., Sakurai, T., Atoda, T., “Crystal Structure of α-AlB12”, J. Solid state Chem., 20, 67–77 (1977), (Experimental, Crys. Structure, 21) Portnoy, K.I., Bogdanov, V.I., Mikhailov, A.V., Fuks. D.L., “Interaction Parameters in Interstitial Solid Solutions Based on Aluminium”, Russ. J. Phys. Chem. (Engl. Transl.), 55(4), 583–584 (1981), translated from Zh. Fiz. Khim., (55), 1041–1043 (1981) (Experimental, 10) Doerner, P., “Constitutional Investigations on High Temperature Ceramics of the B-Al-C-Si-N-O System by Means of Thermochemical Calculations” (in German), Thesis, Uni. Stuttgart, Inst. Metallkunde, (1982) (Calculation, Thermodyn., 126) Higashi, I., “Aluminum Distribution in the Boron Framework of γ-AlB12”, J. Solid State Chem., 47, 333–349 (1983) (Experimental, Crys. Structure, 21) Olesinski, R.W., Abbaschian, G.I., “The B-Si System”, Bull. Alloy Phase Diagrams, 5, 478–484 (1984) (Review, Phase Diagram, 51) Lukas, H.L., “Aluminium - Boron - Silicon”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.20852.1.20, (1990) (Crys. Structure, Phase Diagram, Assessment, 6) Duschanek, H., Rogl, P., “The Al-B (Aluminum-Boron) System”, J. Phase Equilib., 15, 543–552 (1994) (Assessment, Crys. Structure, Phase Relations, Review Thermodyn., 78) Gro¨bner, J., Lukas, H.L., Aldinger, F., “Thermodynamic Calculation of the Al-Si-C System”, Calphad, 20, 247–254 (1996) (Calculation, Phase Diagram, Thermodyn., 37) Fries, S.G., Lukas, H.L., “B-Si” in “COST 507. Thermochemical Database for Light Metal Alloys”, Vol. 2”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Luxembourg, 126–128 (1998) (Calculation, Phase Diagram, Thermodyn., 0) Wang, Li., Bian, Xiufang, Sun, Yimin, “Refining Effect of Boron on Hypoeutectic Al-Si Alloys”, Chin. J. Nonfer. Met., 9(4), 714–718 (1999) (Experimental, Morphology, 13) Meschel, S.V., Kleppa, O.J., “Thermochemistry of Alloys of Transition Metals and Lanthanide Metals with some IIIB and IVB Elements in the Periodic Table”, J. Alloys Compd., 321, 183–200 (2001) (Experimental, Thermodyn., 82) Wang, Li, Bian, Xiufang, “Refining Effect of Boron on Hypoeutectic Al-Si Alloys”, J. Mater. Sci. Technol., 16(5), 517–520 (2000) (Experimental, Morphology, 7) Nogita, K., Dahle, A.K., “Effects of Boron on Eutectic Modification of Hypoeutectic Al-Si Alloys”, Scr. Mater., 48(3), 307–313 (2003) (Kinetics, Morphology, 16) Mirkovic, D., Gro¨bner, J., Schmid-Fetzer, R., Fabrichnaya, O., Lukas, H.L., “Experimental Study and Thermodynamic Re-assessment of the Al-B System”, J. Alloys Comp., 384, 168–174, (2004) (Experimental, Assessment, Calculation, Phase Diagram, Phase Relations, Thermodyn., 26) Groebner, J., Mirkovic, D., Schmid-Fetzer, R., “Thermodynamic Aspects of Grain Refinement of Al-Si Alloys Using Ti and B”, Mater. Sci. Eng. A, 395(1-2), 10–21 (2005) (Assessment, Calculation, Phase Diagram, Phase Relations, Thermodyn., 56) Huiyuan, G., Yanxiang, L., Xiang, C., Xue, W., “Effects of Boron on Eutectic Solidification in Hypoeutectic Al-Si Alloys”, Scr. Mater., 53(1), 69–73 (2005) (Experimental, Morphology, 8) Yoshikawa, T., Morita, K., “Thermodynamic Properties of B in Molten Si and Phase Relations in the Si-Al-B System”, Metal. Trans., 46(6), 1335–1340 (2005) (Experimental, Phase Relations, Thermodyn., 16) Yoshikawa, T., Morita, K., “Removal of B from Si by Solidification Refining with Si-Al Melts”, Metal. Trans. B, 36(6), 731–736 (2005) (Experimental, Phase Relations, 18)

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Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)


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Aluminium – Carbon – Titanium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Lesley Cornish, Gabriele Cacciamani, Damian M. Cupid, Jozefien De Keyzer

Introduction Al-C-Ti alloys are of interest for many different reasons, mainly as alloys in themselves, dispersion-hardened alloys with TiC, and also as grain refiners for commercial Al-based alloys, giving improved properties. Alloys are based on Ti-6Al/TiC [2002Zha2], which has high strength, low density and good elastic modulus. Ti3Al and γαffl (TiAl) phases are being developed for lightweight structural applications, especially at elevated temperatures. The system is also important for composite manufacture; using TiC in (Al), and also using the ternary carbides. Ti2AlC is recognized as a “machinable ceramic”, its layered structure produces a combination of metallic and ceramic properties. Additionally, there is interest for coatings: Ti-Al-C PVD coatings could compete with Ti-Al-N, and in a different application, TiC has potential for a barrier diffusion layer for silicon semiconductor devices. Phase boundary information on this system comes from the work of [1954Thy, 1980Sch, 1987Zak], [1989Cam1, 1989Cam2, 1991Cam1, 1991Cam2, 1992Pie, 1994Pie] mostly using optical and electron microscopy together with and X-ray diffraction. Other publications [1963Jei, 1964Jei, 1964Now, 1976Ivc1, 1976Ivc2, 1980Pea] covered X-ray diffraction studies, crystal structures, lattice parameters and physical properties of the two ternary carbides Ti2AlC and Ti3AlC, with general agreement. One of the first reviews was by [1965Mol], followed by [1983Kub, 1983Sri]. The first evaluation within the ongoing MSIT Evaluation Program was made by [1990Hay], which is updated by the present work. Recent reviews by [1994Pie, 2000Ban, 2006Rag] (where [1994Pie] also undertook experimental work) mainly include the experimental information obtained after the original MSIT review of [1990Hay]: [1990Via, 1991Cam1, 1992Pie, 1994Pie, 1994Zha]. Information of some other recent articles that were not considered there ([1990Hay]) are included in this review [1991Cam2, 1995Via, 2000Ria, 2000Tze, 2000Van, 2003Ge1, 2005Zho2]. An overview of the investigations considering phase equilibria, solid phases and thermodynamics of the system is given in Table 1.

Binary Systems For the Al-Ti system, the MSIT evaluation from [2004Sch] is available. More recently, a new review of [2006Sch] and a thermodynamic assessment of [2007Wit] appeared which show differences in two main regions. [2006Sch] and [2007Wit] consider the equilibria between (αTi), (βTi) and Ti3Al. Starting from the same experimental literature information, [2004Sch] concluded that Ti3Al decomposes congruently and there is no invariant equilibrium involving Landolt‐Bo¨rnstein New Series IV/11E1

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these three phases, while [2006Sch] and [2007Wit] concluded that there are two peritectoid reactions between these three phases and a small two-phase field between (βTi) and Ti3Al in the 1150˚C-1200˚C temperature range. According to [2004Sch], the different phase equilibria may be related to very small differences in Gibbs energy, and consequently in driving forces, to give the possible phase transitions in that range. Thus, it is very difficult to decide which interpretation of the experimental data is actually correct. The second difference considers the equilibria between γ(TiAl), TiAl2 and one-dimensional antiphase domain structures or “long period structures (LP)” stable at high temperature in the range between 65 and 75 at.% Al. In this case, the assessments by [2006Sch] and [2007Wit] are based on more recent experimental data and give a more complete discussion of the stable equilibria in this region. Thus the phase diagram calculated by [2007Wit] is accepted in the present evaluation since it is based on experimental data available, the majority of which were assessed in [2006Sch], being complemented by own key experiments. The adopted Al-Ti phase diagram is shown in Fig. 1 from [2007Wit]. The Al-C phase diagram is accepted from the review of [2004Per]. The C-Ti phase diagram (Fig. 2) is taken from the recent thermodynamic calculation of [2003Fri]. It leads to a better fit of high temperature heat capacity measurements than all previous studies. The phase diagram of [2003Fri] is reproduced with the same accuracy as given by [1999Dum] and gives a better description of the experimental data than previous assessments. However, it should be noticed that the liquidus lines are not supported by any experimental data.

Solid Phases The solid phases are given in Table 2. The different morphologies of TiAl3 are deemed to be related to the different morphologies observed at different cooling rates and Ti contents [1982Bla, 1993Sve]. Similarly, TiAl2 has been reported to have metastable structures in ascast alloys [2001Bra]. The Ti2Al5 phase [2004Sch] is interpreted here as a one dimensional antiphase domain structure or “long period structure (LP)” stable at high temperature in the range between 65 and 75 at.% Al. The Ti5Al11 phase has been included in Table 2, although it is not accepted here as a separate phase, and is presumed to be related to ζ, Ti2Al5, with which it shares similar lattice parameters. According to [1989Cam1, 1989Cam2], carbon solubilities in the binary aluminides are small (~1 at.% C) and the absence of fine scale carbide precipitates after ageing the alloys with 1 at.% C for 15 days at 750˚C and 450˚C respectively shows the absence of a significant decrease of C solubility in (αTi) or Ti3Al [1989Cam1, 1989Cam2]. Ternary carbides are stated to form above 1250˚C by peritectic reactions from TiC1–x and liquid. All workers have reported minimal solubility of aluminium in TiC1–x [1984Ker, 1980Sch, 1989Cam1, 1989Cam2, 1990Cam, 1991Cam1, 1991Cam2, 1994Pie, 2000Ria]. [1980Sch] confirmed the existence of the hexagonal ternary H-phase Ti2AlC and of the cubic Perovskite-(P-) phase Ti3AlC reported by earlier workers [1963Jei, 1964Jei, 1964Now, 1976Ivc1, 1976Ivc2, 1980Pea]. A cubic phase that was earlier noted in a Ti-10Al-1C (mass%) alloy [1954Thy] probably was the cubic Perovskite-phase Ti3AlC; this was confirmed by [2000Ria], although [1994Pie] reported not being able to find a match. [1994Pie] found a third ternary phase, N, Ti3AlC2. It is presumed that the reason for this phase not being observed earlier was that longer sintering times and higher pressures [2006Li] were needed. DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009

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The ternary carbides have less C than their stoichiometric compositions [1980Sch, 1992Pie, 1994Pie, 2000Tze] and the range for N is at least Ti3Al1.1C1.8 [2000Tze] to Ti3Al1.1C2.3 [2005Kho]. The relationship of TiC1–x and Ti2AlC1–x to Al4C3 was discussed by [1982Now]. The unusual diffusion path reported by [1995Via] between (Al) and TiC could possibly have been due to the then unrealized N phase.

Quasibinary Systems The section between γ (TiAl) and TiC1–x by [1987Zak] is not a true quasibinary since both peritectic reactions have a (non-linear) three-phase field.

Invariant Equilibria The invariant reactions are not fully derived yet, and more work needs to be done. The agreed invariant equilibria of [1991Cam1, 1994Pie] are given in Table 3. Ti2AlC melts incongruently at 1625±10˚C; Ti3AlC also melts incongruently at 1580±10˚C, whereas Ti3AlC2 decomposes in the solid state [1994Pie]. [2004Hwa] observed that ternary carbides formed from peritectic reaction with TiC1–x ; this confirmed the reaction: L + TiC1–x Ð H of [1987Zak], which [1994Pie] reported having a maximum at 1625±10˚C. [1992Che] deduced that cubic Ti3AlC is a metastable transition phase which is ultimately replaced by the more stable H (Ti2AlC) phase, which agrees with the reactions of [1994Pie]. According to [2005Zho2], the lower stability limit of Ti3AlC2 is ~1300˚C (decomposing to give TiC1–x and Al vapor) and [2000Tze] reports the upper stability limit as at least 1400˚C, if not 1450˚C. The most Al rich equilibrium in Table 3 is from [1990Via, 1993Sve], where the latter deduced that the formation kinetics of Al4C3 were slow, and the reaction agreed with [1998Fra]. [1990Via] gave alternative Al rich reactions: L + Al4C3 + TiAl3 Ð (Al) or L + TiAl3 Ð Al4C3+ (Al), both being very near to the melting point of aluminium.

Liquidus, Solidus and Solvus Surfaces [1991Cam1] derived a “semi-schematic” liquidus surface. The volume fraction of H was less than the projection suggested; this was thought to be due Al substituting for C. A schematic liquidus surface and partial reaction scheme was also derived by [1994Pie]. They are not presented here, since they contain extra Al-Ti phases which are not accepted in this review. However, the liquidus surface is dominated by TiC1–x , then graphite. The reactions forming the other phases all lie very close to the Al-Ti binary. [1990Via] gave an experimental liquidus projection in the Al rich corner, whereas [1998Fra] produced one from thermodynamic calculations. The major difference was that [1990Via] gave the lowest temperature reaction, which was dependent on cooling conditions. [1993Sve] also attempted to rationalize the different morphologies of TiAl3 observed with composition.


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Isothermal Sections [2003Ge1] drew an isothermal section at 1300˚C showing the third ternary phase, N (Ti3AlC2), which was involved in the three-phase triangles: (N + H + TiC1–x), (N + TiC1–x + TiAl3) and (N + H + Ti5Al11), but the last phase (as well as its Al-Ti phase diagram) is not accepted here. Additionally, the phase compositions did not agree with the boundaries due to Al loss. Similarly, the isothermal section at 1300˚C of [1992Pie, 1994Pie] is not presented. The 1100˚C isothermal section of [1994Zha] is not presented because it contains Ti2Al5 which is not accepted here, although it mainly agrees with the other sections presented. [2000Ban] gave isothermal sections at 1300, 1100 and 1000˚C, but these are not reproduced because they are incompatible with the Al-Ti binary accepted here, and show a radical change in the positions of the P and H ternary phases over the temperature ranges (especially 1100˚C). [1993Tia] established a metastable diagram between 900 and 1050˚C. The isothermal section at 1250˚C was drawn by [1991Cam1], and is shown in Fig. 3 after modifications to be consistent with the accepted Al-Ti binary. [2000Ria] modified the 1050˚C isothermal section of [1990Cam], with additional results that retained the same phase fields, but changed the shape of some of them, notably ((αTi) + P + TiC1–x). This is redrawn in Fig. 4 to be consistent with the Al-Ti binary system accepted here, the positions of the P and H, and keeping their alloys in the respective phase fields. The redrawn P and H phases in Figs. 4 and 1 are also consistent with those of [1990Via], but the isothermal sections of [1990Via] are not given here because negligible solubility is given for all binary phases. The results of [1954Thy] and [1989Cam1, 1989Cam2, 1991Cam1] indicate that both (αTi) and (βTi) solid solutions are in equilibrium with TiC, thus generating the tie-triangles ((αTi) + (βTi) + TiC) and ((αTi) + TiC + P) (Fig. 5) at 1000˚C. This is taken rather than the Ti rich corner of [1980Sch], which showed that at 1000˚C, (βTi) solid solution is in equilibrium with the Perovskite carbide phase Ti3AlC (P) generating the two tie-triangles ((βTi) + TiC + P) and ((αTi) + (βTi) + P). This decision was made because [1980Sch] showed those relationships as dotted, and did not give the sample compositions. Their drawn phase relationships also contradict their statement that (βTi) was not retained in the ternary. Additionally, seven experimental alloys of [1994Zha] agree with the findings and interpretation of all previous workers, except for [1980Sch]. There is agreement that γαffl (TiAl) is in equilibrium with Ti2AlC at 1000˚C [1980Sch, 1987Zak, 1989Cam1, 1991Cam1]. The isothermal section at 750˚C (Fig. 6) is taken from [2000Ria], but redrawn to be consistent with the Al-Ti binary accepted here (Fig. 1), the positions of the P and H, as well as the low carbon solubilities reported in the Al-Ti phases, except for (αTi) [1954Thy, 1980Sch, 1987Zak, 1989Cam1, 1989Cam2, 1991Cam1, 1991Cam2, 1992Pie, 1994Zha]. This is consistent with the isothermal section at 1000˚C in terms of reduced ternary element solubilities in the phases at lower temperatures. The section of [2000Ria] is preferred to that of [1991Cam1] because it was derived using alloys further into the ternary system (up to 15 at.% C) [2000Ria], instead of only up to 3 at.% C [1991Cam1]. The redrawing of the diagram brings the only contentious alloy of [1991Cam1] to a position nearly on the boundary between the ((αTi) + TiAl3 + P) and ((αTi) + P + TiC1–x), which could explain why P was not found. Although it appears that TiC1–x [1991Cam1] was re-precipitated within the grains, it would have been the primary phase [1994Pie] and [1995Via] has shown that it takes a long time to disappear. In the isothermal sections presented, the compositions of the P and H phases were modified to agree with those of the 1000˚C isothermal section since this was the widest ranges, and DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009

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otherwise the other sections had compositions which fell outside this range, especially those of [2000Ria] which used higher Ti content alloys. Experimental work on liquid (Al) by [1993Sve] showed that there were three-phase regions: ((Al) + TiAl3 + TiC1–x) and ((Al) + Al4C3 + TiC1–x). Several calculations have been undertaken on the Al rich corner: [1998Fra] showing the equilibrium between (Al) with TiC1–x and TiAl3 and [2000Van] which confirmed [1998Fra], while studying grain refiners.

Temperature – Composition Sections Isopleths for the concentration range 0–1 mass% C at 2, 4, 6, 8 and 10 mass% Al were determined by [1954Thy] and boundaries between the following phase fields were located: (αTi), (βTi), ((αTi) + (βTi)), ((βTi) + TiC1–x), ((αTi) + (βTi) + TiC1–x) and a field containing more than one phase one of which is TiC1–x. The solubility of C in (αTi) was found to increase with increasing Al contents rising to over 1 mass% C at 1150˚C and 10 mass% Al. These are reproduced in Figs. 7–11) from [1990Hay] who redrew them slightly (but without compromising the original experimental results) to be consistent with isothermal sections, and they are redrawn slightly to be consistent with the adopted Al-Ti binary. A section at 0.2 at.% C [1997Li, 1999Li] agrees with the accepted binary. [1987Zak] derived a γ(TiAl)TiC1–x section (Fig. 12).

Thermodynamics [1986Ban] determined the Gibbs free energy of formation of the compounds Al4C3 and TiC as a function of temperature in the range 300–1800 K. The same was calculated by [1991Rap] in order to demonstrate that the formation of Al4C3 is metastable with respect to TiC and it is due to kinetic reasons. Thermodynamics and kinetics of formation of TiC by a reaction between Al-Ti melts and a carbonaceous gas have been studied by [1991Sah]. More recently [2006Ton] used thermodynamic analysis to show that it is feasible to produce TiC-Al (Ti) nanocomposite powders using Al and Ti as starting materials and CH4 as reacting gas in a thermal plasma environment. Based on thermodynamic results previously available in literature, [1991Yok] presented a three-dimensional potential phase diagram for the Al-C-Ti system at 700˚C. [1993Jar] investigated thermodynamics and kinetics of the reactions taking place during infiltration of graphite fibres by molten Al-Ti alloys. A temperature range for safe processing of Al-TiC composites has been suggested by [1993Mit] based on thermodynamic considerations. [1993Sve] determined by a theoretical thermodynamic analysis the C and Ti concentration in Al-rich liquid in equilibrium with solid phases at temperatures between 700 and 1100˚C. A thermodynamic discussion of the solid/liquid phase equilibria in the Al rich corner has been reported also by [1998Fra]. Limiting values of thermochemical data of formation of the Al-C-Ti ternary compounds are reported in [1994Pie], based on binary and ternary Al-C-Ti equilibria and thermochemical data for various reactions previously studied in literature. The authors concluded that the dataset so obtained was sufficient to calculate tie lines in agreement with known phase equilibria but a full thermodynamic assessment of the system is needed. Landolt‐Bo¨rnstein New Series IV/11E1

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Thermodynamics of formation of TiC in DC-casting and melt spun Al-C-Ti alloys has been discussed by [1994Ton]. Thermodynamics and kinetics of formation of TiC in Al-C-Ti have also been investigated by [2000Zha2]. Thermodynamics and kinetics of precipitation of TiC in Al-C-Ti melts at 1100˚C have been investigated by [1999Jar]. Low temperature heat capacity of Ti3Al1.1C1.8 has been measured by [1999Ho]. Heat capacity, thermal expansion coefficients and thermal conductivity of the Ti2AlC phase have been measured by [2002Bar] at room to high temperature: these are reported in Table 4. Low temperature heat capacity of the same phase has been measured by [2006Dru].

Notes on Materials Properties and Applications The work on materials properties is summarized in Table 5. The system is of interest because of the effect of different additions to the grain refining of many commercial Al-based alloys [1949Cib], with the benefits of improved mechanical properties, reduced ingot cracking, improved feeding and increased casting speeds, improved homogeneity and reduced porosity and better deformation behavior. Grain refiners based on Al-C-Ti can be added in master alloy form, and are an alternative to conventional Al-B-Ti grain refiners [1994May] and are less susceptible to agglomeration. [1976Jon] calculated solubility limits of liquid aluminium in different systems, including Al-C-Ti. Epitaxy was found to be important; thus TiC was seen as being more effective than TiAl3, having more epitaxial faces [1972Cis]. However, TiC is surrounded by Al4C3, then Ti3AlC, which reduces the grain refinement effectiveness, and so heating liberates the beneficial TiC [1986Ban]. Several studies were made on the formation of the grain refining precipitates [1977Ivc, 1989Cam1, 1991Fro, 1991Sah, 2000Bri1, 2000Zha1, 2002Zha1, 2003Zha2], to improve the methods [2000Zha1, 2005Zha, 2005Kum] and to assess the efficiency [2000Bri2, 2000Rom, 2000Sch, 2000Van, 2002Gaz, 2002Tro, 2002Xia, 2002Liu, 2003Zha2]. Aluminium alloys grain refinement has also been achieved by synthesizing Al-Ti-C launders and moulds in an electromagnetic field [2005Xia]. There are also several alloys based on Ti-6Al/TiC1–x [2002Zha2] which have high strength, low density and good elastic modulus. Aluminium dispersion-hardened alloys are also important and different alloys in different conditions have been studied: up to 50 mass% TiC1–x [1986Bat]. The Ti3Al and γ (TiAl) phases are being developed for lightweight structural applications, especially at elevated temperatures. γ (TiAl) is disadvantaged by low temperature ductility and toughness, and these are being improved by precipitation [1991Mab]. Composites are also being developed by precipitating Ti3AlC in HIPped γαffl (TiAl) alloys [1992Che, 1994Whi], although the precipitate long-term stability was questioned [1992Che]. Optimization of the lamellar microstructural has been attempted [2001Qin]. Improvements have also been made by indirect-extruding TiAl-Mn-Mo-C alloys to obtain a fine microstructure, and rheocasting TiAl alloys and composites [1997Ich]. The system is also important for composite manufacture; composites have been made by infiltrating graphite fibres with molten Al-Ti, using the phase diagram [1993Jar]. In-situ internal carburization has also been carried out to fabricate the composites [1991Jar, 1999Bir, 2001Zha]. [1995Mor] made in-situ TiC1–x particulate reinforced aluminum composites by reacting graphite particles in a liquid Al-Ti alloy and showed that the Young’s modulus and tensile strength increased proportionally to the TiC1–x volume fraction. Ductile Al/TiC DOI: 10.1007/978-3-540-88053-0_6 ß Springer 2009

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MMCs (metal matrix composites) have also been made [1993Mit, 2003Arp] and the effects of the TiC1–x dendrite spacing were studied [2002Ria]. Bulk [2004Hwa, 2004Wan3, 2005Jia, 2005Li, 2006Xia] and porous [1992Vol] Al-C-Ti cermets have also been manufactured using the SHS (self-propagating high-temperature synthesis) sintering technique [1992Vol, 1997Lee, 1997Ye, 2002Mei, 2006Xia] and phase evolution studied [2005Rog]. Variations of thermal explosion/quick pressure (TE/QP) methods have also been used [2001Yan, 2005Ma], and “hot shock” [1998Tan]. Combustion synthesis, also known as self-propagating high-temperature synthesis (SHS), offers advantages of low processing cost, energy-efficiency, and high production rate, and has been used to manufacture TiC1–x [1993Cho, 2006Xia]. Addition of synthesized master alloys to molten aluminium has also been achieved [2006Sel], and a dipping exothermic reaction process (DERP) has also been used [2003Son, 2004Son2]. [1996Tom] produced ternary composites by combustion synthesis with titanium, aluminium, and graphite powders. At low C-contents, binary intermetallics were formed, whereas Ti2AlC and TiC1–x were formed at higher C-contents. Other elements have also been added to the materials [2003Yan]. A useful review of these and other composites is [2000Tjo]. Some of the composites are being developed for cutting tools [1967Hod, 1976Ivc1], using the intermetallic phases, and TiC1–x shows the most promise. An extension of this are complex carbides (or nitrides) M2AlC (or N), where M = Cr, Ti, V, Nb or Ta [1996Cha]. Also, the system has been considered for coatings on Ti-based metal matrix composites containing alumina fibres [1994Din]. Studies have been made of solid state reactions at graphite/TiAl or Ti3Al interfaces, showing that whilst graphite moulds are safe for these alloys, carbon fibres would have to be coated to prevent Ti diffusion [1995Via]. Ti2AlC is recognized as a “machinable ceramic” (layered structure) with good mechanical properties and low thermal expansion, thus giving potential high temperature applications (furnace apparatus). It has been manufactured by self-propagating high-temperature synthesis (SHS) [2001Lop], combustion synthesis [2004Kex], “thermal explosion” (a form of SHS) [2002Kho1, 2002Kho2], hot pressing [2004Hon, 2005Pen], spark plasma sintering [2003Zho, 2004Mei, 2005Zho1], two-stage pressureless sintering followed by hot pressing [2005Wan1]. Similarly, interest has also been shown in Ti3AlC [2001Lop] and Ti3AlC2 [2000Tze, 2001Lop, 2003Ge1, 2003Ge2, 2004Son1, 2005Kho, 2005Zho2, 2006Li], with similar manufacturing processes. Ti3AlC pins twins in γ (TiAl) [1997Chr]. Ti-Al-C PVD coatings have been manufactured as possible competitors to Al-N-Ti coatings [1991Kno] and for coatings for steels [2005Wal]. Another coating application is as a diffusion barrier for materials in contact structures for silicon semiconductor devices. Aluminium can be used (rather than the more effective and expensive PtSi), but needs a barrier diffusion layer, and this can be TiC1–x [1982Wit].

Miscellaneous The formation mechanism for TiC1–x from the ball-milled elemental components was studied by [2003Wan2]. Various reaction mechanisms for combustion reactions to yield TiC1–x [1993Cho] (with results of the first TiC formation supported by [1995Via]), Ti3AlC2 [2002Ge], and TiC / γ(TiAl) composites [2002Hwa] have been proposed. [1993Cho] studied the effect of Al addition (0-40 mass%) on the combustion reaction between Ti and C to form TiC1–x using combustion wave velocity data and differential thermal analysis. Their results Landolt‐Bo¨rnstein New Series IV/11E1

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indicate the reaction between titanium and aluminium initiates the reaction between titanium and carbon through a series of reaction steps. [2002Hwa] also used combustion synthesis to create TiC1–x/Al composites, but reported the presence of ternary carbides in the final microstructure. In-situ processing of aluminium matrix metal matrix composites based on TiC has been investigated [1995Mor, 1996Tom, 1993Ber]. [1996Mat] performed electronic structure calculations for C-containing γ (TiAl) and showed that C substitutes Al sites. [1996Ily] undertook electronic structure calculations on Ti1–xAlxC (x = 0.17), and showed that this compound has a bandgap of 2.18 eV. Ab-initio electronic structure calculations were performed on C-containing γ (TiAl) [1996Mat] and on theoretical Ti1–xAlxC compounds (NaCl lattice where Al and Ti occupy the same lattice sites) by [1996Ily, 1997Ily1, 1997Ily2, 1998Mat]. The wetting angle of molten Al on TiC1–x was measured using the sessile drop method by [1966Yas] at 117˚. [1996Sob] used the same method to investigate the wettability of Al-rich Al-Ti alloys (Al-0.5Ti to Al-10Ti (mass%)) and showed that Ti additions lead to better wettability of graphite because of reactions between titanium and carbon. Gibbs free energy calculations [1994Ton, 1996Tom, 1995Mor] have been used to predict/ determine reactions and phase stabilities for aluminium-based metal matrix composites with Ti and C. [1997Wu] ball milled titanium, aluminium, and graphite powders in the atomic ratio Al-0.23Ti-0.23C and produced an amorphous phase that further transformed to an fcc metastable phase. Annealing at 600˚C for 45 minutes yielded the equilibrium phases (Al), Al3Ti and TiC1–x(?).

. Table 1 Investigations of the Al-C-Ti Phase Relations, Structures and Thermodynamics

Reference


Temperature/ Composition/Phase Range Studied

[1954Thy]

Optical metallography and XRD

0-1 mass% C at 2, 4, 6, 8 and 10 mass% Al

[1963Jei]

Powder metallurgy, XRD

Ti2AlC

[1964Jei]


Ti3AlC

[1964Now]


Ti2AlC

[1972Now]


Ti2AlC

[1976Ivc2]


Ti2AlC and Ti3AlC

[1977Gur]

Metallography, XRD, magnetic properties TiAl2.97C0.03 measurements

[1980Pea]

Dimensional analysis

Ti2AlC

[1980Sch]

XRD

1000˚C; 30 alloys of nonspecified composition

[1984Ker]

Metallography, SEM and EDS.

Solubility of Al in TiC1–x at 1600˚C

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Reference

Temperature/ Composition/Phase Range Studied


[1987Zak]

Metallography and X-ray diffraction

600˚C - 1400˚C, TiAl-TiC, up to 25 mass% TiC0.98

[1993Tia]

XRD, TEM

(Ti0.50Al0.50)99.5C0.5

[1989Cam1, 1989Cam2, 1990Cam, 1991Cam1, 1991Cam2]

Light microscopy, SEM, TEM, X-ray diffraction, X-ray ED spectroscopy, EDX and hardness measurements

750˚C, 1000˚C, 1250˚C, 15 < at.% Al < 55, 0.5 < at.% C 1 50 at.% Al at 1000˚C [2001Bra] 62 at.% Al at 1100˚C [2001Bra]

Ti1–xAl1+x > 1170

tP4 P4/mmm AuCu

a = 403.0 c = 395.5

Ti3Al5 < 809

tP32 P4/mbm Ti3Ga5

a = 1129.3 Ti3Al5, stable below 810˚C [2001Bra] (a = 399.3 substructure) c = 403.8

ζ, Ti2Al5 1432 - 976

tP28 P4/mmm Ti2Al5

a = 393.6 c = 413.54

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Disordered Ti1–xAl1+x has c/a < 1 [2001Bra], and was reported as a separate high temperature phase

[1986Mii] Presumed ordering effect

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Lattice Parameters [pm] a = 398.0 c = 2436.0

Comments/References

TiAl2 < 1224

tI24 I41/amd HfGa2

Ti5Al11 ≥ 995

CuAu a = 395.3 superstructure c = 410.4

[2001Bra] 1300˚C, but not recognized here as a separate phase

ε(h), TiAl3 (h) 1396 - 734

tI8 I4/mmm TiAl3(h)

[V-C2]

a = 384.6 c = 859.4

[1980Mii, 1986Mii] by TEM

a = 384.9 1000˚C [2001Bra] c = 860.9 (c = 430.5 substructure)

ε(l), TiAl3 (l) < 932

tI32 I4/mmm TiAl3(l)

a = 387.7 76 at.% Al at 640˚C [2001Bra] c = 3382.82 (c = 422.9 substructure)

TiC1–x

cF8 Fm3m NaCl

a = 432.8

[V-C2]

Al4C3

hR21 R3m Al4C3

a = 333.8 c = 2511.7

[V-C2]

* H, Ti2AlC1–x

hP8 P63/mmc Cr2AlC

a = 305.6 c = 1362.3

[V-C2]

* P, Ti3AlC1–x

cP5 Pm3m CaTiO3

a = 415.6

[V-C2]

* N, Ti3AlC2

hP12 P63/mmc CMo

a = 307.53 c = 1857.8

[1994Pie], isotypic with Ti3SiC2

Metastable / high pressure phases TiAl3 (m)

cP4 Pm3m AuCu3

Ti3Al (I) hP16 15 to > 41 GPa P63/mmc TiNi3


a = 397.2

Splat cooled; 85 at.% Al [1994Bra]

a = 531.2 c = 960.4

[1997Sah] at 16 GPa; Not found at 0–35 GPa, 25-2250˚C [2004Sch]

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Comments/References

TiAl2 (m)

oC12 Cmmm ZrGa2

a = 394.21 c = 401.61

~66 to 67 at.% Al - metastable in as-cast alloys [2001Bra]

Ti52Al48 (m)

cP20 P4132 βMn

a = 690

Precipitated in amorphous thin film after 1 h at 525˚C [1999Abe]

. Table 3 Invariant Equilibria Composition (at.%) T [˚C]

Reaction

Type

Phase

Al

C

Ti

L + TiC1–x + H Ð P

1580±10

P

-

-

-

-

L + TiC1–x Ð (βTi) + P

1430 < T < 1580

U

-

-

-

-

L + P Ð (βTi) + H

Lower than the above

U

-

-

-

-

L + (βTi) Ð (αTi) + H

776˚C [V-C2]

x = 0, 0 < y ≤ 0.047,T = 1105˚C [1990Kas] y = 0, 0 < x ≤ 0.42, T = 1284˚C [Mas2] in the U95.2Mo2.8Al2.0 and U93.9Mo2.1Al4.0 (at.%) alloys annealed at 900˚C 48 h and cooled to room temperature at 4–5˚C·min–1, together with α and λ2 phases [1964Nic] in the U87.5Mo12.0Al0.5 (at.%) alloy homogenized at 1050˚C, reduced and quenched from 800˚C [1984Gom] T > 668˚C [V-C2]

x = 0, 0 < y ≤ 0.0054, T = 758˚C [1990Kas] y = 0, 0 < x ≤ 0.02, T = 668˚C [Mas2] T = 25˚C [V-C2]

x = 0, 0 < y < 7·10–4, T = 665˚C [1990Kas] y = 0, 0 < x ≤ 0.02, T = 570˚C [Mas2] 0 ≤ x ≤ 0.27 [1971Pet]

a = 778

[V-C2]

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Pearson Symbol/ Lattice Space Group/ Parameters Prototype [pm]

Comments/References

UAl3 < 1350

cP4 Pm3m AuCu3

a = 426.2

UAl4 < 731

oI20 Imma UAl4

a = 440.14 [2004Tou] b = 625.52 c = 1372.79

Mo3Al ≲ 2150

cP8 Pm3m Cr3Si

ξ2, MoAl 1750 - 1470

[V-C2]

a = 495

23 to 28.5 at.% Al at T ≈ 1720˚C [2005Sch] [1958Woo, 2005Sch]

a = 309.8

46 at.% Al at T ≈ 1720˚C to 52 at.% Al at T = 1570˚C [2005Sch] [1971Rex, 2005Sch]

cP2 Pm3m CsCl

ξ1, Mo2Al3 1570 - 1490

-

-

63 at.% Al [2005Sch]

Mo3Al8 < 1555

mC22 C2/m Mo3Al8

a = 916.4 b = 363.9 c = 1004 β = 100.5˚

[V-C2]

MoAl3 1222 - 818

mC32 C2/m MoAl3

a = 1639.6 [2005Sch] b = 359.4 c = 838.6 β = 101.88˚

Mo1–xAl3+x 1260 - 1154

cP8 Pm3m Cr3Si

76 at.% Al at T = 1222˚C to 78 at.% Al at T = 1154˚C [2005Sch] a = 494.5

[2005Sch]

MoAl4 1177 - 942

mC30 Cm WAl4

Mo4Al17 < 1034

mC84 C2 Mo4Al17

a = 915.8 b = 493.23 c = 2893.5 β = 96.71˚

[1995Gri, 2005Sch]

Mo5Al22 964 - 831

oF216 Fdd2 Mo5Al22

a = 7382 b = 916.1 c = 493.2

[1995Gri, 2005Sch]


79 to 80 at.% Al at T = 1154˚C [2005Sch] a = 525.5 [V-C2] b = 1776.8 c = 522.5 β = 100.88˚

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Pearson Symbol/ Lattice Space Group/ Parameters Prototype [pm]

MoAl5 (h2) 750 - 800 < T < 846

hP12 P6322 WAl5

a = 491.2 c = 886.0 a = 493.7 c = 924.3

[2005Sch]

MoAl5 (h1) 648 < T < 750–800

hP60 P321

a = 493.3 c = 4398

[2005Sch]

Comments/References

Space group [P63V-C2]

MoAl5 (h1) MoAl5 (r) ≲ 648

hP36 R3c MoAl5 (r)

a = 495.1 c = 2623

[2005Sch]

MoAl12 < 712

cI26 Im3 WAl12

a = 757.3

[1954Ada, 2005Sch]

U2Mo ≲ 1252 (?)

tI6 I4/mmm MoSi2

a = 342.7 c = 983.4

* τ1, U9Mo16Al75

-

-

[1969Pet]

-

0.31 ≤ x ≤ 0.40 [1969Pet]

* τ2, U(MoxAl1–x)2 hP12 P63/mmc MgZn2

32.5 to 34 at.% Mo [Mas2] [V-C2]

* τ3, U6Mo4Al43

hP2 a = 1096.6 P63/mcm c = 1769.0 Ho6Mo4+xAl43–x (x = 0.11) or Yb6Cr4+xAl43–x (x = 1.15)

[1994Nie]

* τ4, UMo2Al20

cF8 Fd3m CeTi2Al20 or CeCr2Al20

[1995Nie]

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a = 1450.6

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T [˚C]

Type

Phase

L + ξ1 Ð Mo3Al8 + ξ2

1510

U1

L

L + ξ2 Ð Mo3Al8 + Mo3Al

L + λ2 Ð τ2

L Ð τ2 + Mo3Al

1480

1460

1410

U2

p3

e5

1380

U3

ξ1

63.0

37.0

0

Mo3Al8

72.7

27.3

0

ξ2

48.0

52.0

0

60.0

30.0

ξ2

48.0

52.0

0

Mo3Al8

72.7

27.3

0

Mo3Al

25.0

75.0

0

L

45.2

21.5

33.3

L

L Ð λ2 + Mo3Al8 + Mo3Al

1340

U4

E1

1220

U5

48.2

18.5

33.3

46.0

20.7

33.3

L

39.0

29.0

32.0

τ2

39.7

27.0

33.3

25.0

75.0

L

48.5

27.5

24.0

τ2

45.5

21.5

33.3

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0

49.0

18.0

25.0

75.0

66.4

18.3

15.3

τ1

75.0

16.0

9.0

λ2

61.0

6.0

33.0

Mo3Al8

72.7

27.3

0

L

58.0

26.0

16.0

58.5

8.5

33.0

Mo3Al8

72.7

27.3

0

Mo3Al

25.0

75.0

0

L

6.0

28.0

66.0

(Mo)

1.5

97.0

1.5

γ

1.0

39.0

60.0

25.0

75.0

L

Mo3Al


10.0

λ2

λ2

L + (Mo) Ð γ + Mo3Al

4.0

τ2

Mo3Al 1352

U

29.5

λ2 L + τ1 Ð Ð λ2 + Mo3Al8

Mo 66.5

Mo3Al L + τ2 Ð λ2 + Mo3Al

Al

33.3 0

0

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Al–Mo–U


T [˚C]

Type

Phase

L + Mo3Al Ð γ + τ2

1140

U6

L

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U7

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Mo

U

9.0

22.0

25.0

75.0

γ

3.0

24.0

73.0

τ2

39.2

27.0

33.8

L

11.5

10.5

78.0

Mo3Al

L + τ2 Ð γ + λ 2

Al

69.0 0

τ2

44.8

21.5

33.7

γ

4.5

11.5

84.0

λ2

50.4

15.8

33.8


10

. Fig. 1a Al-Mo-U. Partial reaction scheme, part 1

Al–Mo–U


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. Fig. 1b Al-Mo-U. Partial reaction scheme, part 2

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. Fig. 2 Al-Mo-U. Partial liquidus surface projection


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Al–Mo–U

. Fig. 3 Al-Mo-U. Partial isothermal section at 1250˚C (U-UAl2-Mo3Al8-Mo region)

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Al–Mo–U


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. Fig. 6 Al-Mo-U. Temperature - composition section UAl2-Mo3Al8


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Al–Mo–U

. Fig. 7 Al-Mo-U. Isopleth at 33.3 at.% U

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. Fig. 8 Al-Mo-U. Temperature - composition section U52Al48-Mo3Al (plotted in at.%)


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Al–Mo–U

. Fig. 9 Al-Mo-U. Isopleth at 40 at.% Mo

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. Fig. 10 Al-Mo-U. Isopleth at 20 at.% Al


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Al–Mo–U

. Fig. 11 Al-Mo-U. Isopleth at 50 at.% Al

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. Fig. 12 Al-Mo-U. Isopleth at 20 at.% Mo


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. Fig. 13 Al-Mo-U. Isopleth at 80 at.% U

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References [1954Ada] [1958Woo]

[1964Nic]

[1969Pet] [1971Pet] [1971Rex]

[1979Sof]

[1980Fer] [1984Gom]

[1990Kas]

[1993Ale]

[1994Nie] [1995Gri] [1995Nie]

[1997Lee]

[2000Han]

[2002Kim]

[2002Lee]

[2002Mey]

Adam, J., Rich, J.B., “The Crystal Structure of WAl12, MoAl12 and (Mn, Cr)Al12”, Acta Cryst., (7), 813–816 (1954) (Crys. Structure, Experimental, 14) as quoted by [2005Sch] Wood, E.A., Compton, V.B., Matthias, B.T., Corenzwit, E., “β-Wolfram Structure of Compounds Between Transition Elements, Gallium and Antimony”, Acta Crystallogr., (11), 604–606 (1958) (Crys. Structure, Experimental, 13) as quoted by [2005Sch] Nicholson, S., Harris, D.G., Stobo, J.J., “The Effect of Ternary Additions on the Microstructures of Dilute U-Mo Alloys”, J. Nucl. Mater., 12(2), 173–183 (1964) (Crys. Structure, Morphology, Phase Relations, Experimental, 13) Petzow, G., Rexer, J., “Liquid Equilibria in the Uranium-UAl2-Al8Mo3-Mo System” (in German), Z. Metallkd., 60(5), 449–453 (1969) (Morphology, Phase Diagram, Phase Relations, Experimental, #, 14) Petzow, G., Rexer, J., “Phase Equilibria in the Uranium-UAl2-Al8Mo3-Mo System” (in German), Z. Metallkd., 62(1), 34–38 (1971) (Phase Diagram, Phase Relations, Experimental, #, 9) Rexer, J., “Phase Equilibria in the System Al-Mo at Temperatures above 1400˚C” (in German), Z. Metallkd., 62, 844–848 (1971) (Crys. Structure, Phase Diagram, Experimental, 23) as quoted by [2005Sch] Sofronova, R.M., Nikolaev, A.G., Lyutina, E.M., Voytekhova, E.A., “Influence of Al, Mn, Pd, Ir and Pt Additions on Martensitic Transformation γ → α’b in Uranium Alloys” (in Russian), Alloys for Atomic Energy, Ivanov, O.S., Alekseeva, Z.M. (Eds.), Nauka, Moscow, 131–134 (1979) (Crys. Structure, Morphology, Phase Relations, Experimental, 7) Ferro, R., Marazza, R., “Crystal Structure and Density Data”, Atomic Energy Rev.: Spec. Iss., (7), 359–507 (1980) (Crys. Structure, Phase Relations, Review, 961) Gomozov, L.I., Pokrovskii, A.A., “Influence of Additional Alloying on α-Phase Transformation Kinetics in a Uranium Alloy with 12 at.% Molybdenum”, Russ. Metall. (Engl. Transl.), (4), 142–146 (1984), translated from Izv. Akad. Nauk SSSR, Met., (4), 136–140 (1984) (Crys. Structure, Morphology, Phase Relations, Experimental, Kinetics, Mechan. Prop., 11) Kassner, M.E., Adamson, M.G., Adler, P.H., Peterson, D.E., “The Al-U (Aluminium-Uranium) System”, Bull. Alloy Phase Diagrams, 11(1), 82–89 (1990) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, 44) Alekseeva, Z.M., “Aluminium - Molybdenum - Uranium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.22243.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, #, 6) Niemann, S., Jeitschko, W., “Ternary Aluminides A6T4Al43 with A = Y, Nd, Sm, Gd-Lu, Th, U and T = Cr, Mo, W”, Z. Metallkd., 85(5), 345–349 (1994) (Crys. Structure, Experimental, 23) Grin, Y.N., Ellner, M., Peters, K., Schuster, J.C., “The Crystal Structure of Mo4Al17 and Mo5Al22”, Z. Krist., 210, 96–99 (1995) (Crys. Structure, Experimental, 11) as quoted by [2005Sch] Niemann, S., Jeitschko, W., “Ternary Aluminides AT2Al20 (A = Rare Earth Elements and Uranium; T = Ti, Nb, Ta, Mo, and W) with CeCr2Al20-Type Structure”, J. Solid State Chem., 114, 337–341 (1995) (Crys. Structure, Experimental, 18) Lee, D.B., Kim, K.H., Kim, C.K., “Thermal Compatibility Studies of Unirradiated U-Mo Alloys Dispersed in Aluminium, J. Nucl. Mater., 250(1), 79–82 (1997) (Morphology, Experimental, Interface Phenomena) cited from abstract Han, Y.-S., Park, J.-M., Kim, K.-H., Lee, Y.-S., Kim., C.-K., “An Investigation on the Green Properties of U-10 wt% Mo/Al and U3Si2/Al Powder Compacts”, Nucl. Eng. Design, 202(1), 1–9 (2000) (Morphology, Experimental, Mechan. Prop., 24) cited from abstract Kim, K.H., Park, J.M., Kim, C.K., Hofman, G.L., Meyer, M.K., “Irradiation Behavior of Atomized U-10 wt.% Mo Alloy Aluminum Matrix Dispersion Fuel Meat at Low Temperature”, Nucl. Eng. Design, 211 (2-3), 229–235 (2002) (Morphology, Experimental, Kinetics) cited from abstract Lee, J.-S., Lee, Ch.-H., Kim, K.H., Em, V., “Study of Decomposition and Reactions with Aluminum Matrix of Dispersed Atomized U-10 wt% Mo Alloy”, J. Nucl. Mater., 306(2-3), 147–152 (2002) (Crys. Structure, Morphology, Phase Relations, Experimental, Kinetics, 6) Meyer, M.K., Hofman, G.L., Hayes, S.L., Clark, C.R., Wiencek, T.C., Snelgrove, J.L., Strain, R.V., Kim, K.H., “Low-Temperature Irradiation Behavior of Uranium-Molybdenum Alloy Dispersion Fuel, J. Nucl. Mater., 304(2-3), 221–236 (2002) (Morphology, Experimental, Kinetics) cited from abstract


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[2003Lee]

[2003Mir]

[2003Ryu]

[2004Kei]

[2004Lee]

[2004Tou] [2005Sch]

[2006Kim]

[2006Ryu1] [2006Ryu2]

[2006Wie]

[2007Noe]

[2007Sob]

[Mas2] [V-C2]

Al–Mo–U Garces, J.E., Marino, A.C., Bozzolo, G., “Theoretical Description of the Interdiffusion of Al in the U-Mo Solid Solution”, Appl. Surf. Sci., 219, 47–55 (2003) (Morphology, Calculation, Theory, Interface Phenomena, 11) Lee, S.H., Kim, J.C., Park, J.M., Kim, C.K., Kim, S.W., “Effect of Heat Treatment on Thermal Conductivity of U-Mo/Al Alloy Dispersion Fuel”, Intern. J. Thermophys., 24(5 Special Issue SI), 1355–1371 (2003) (Morphology, Calculation, Experimental, Phys. Prop., 10) cited from abstract Mirandou, M.I., Balart, S.N., Ortiz, M., Granovsky, M.S., “Characterization of the Reaction Layer in U-7 wt% Mo/Al Diffusion Couples”, J. Nucl. Mater., 323(1), 29–35 (2003) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, Kinetics, Transport Phenomena, 18) Ryu, H.J., Han, Y.S., Park, J.M., Park, S.D., Kim, C.K., “Reaction Layer Growth and Reaction Heat of U-Mo/Al Dispersion Fuels Using Centrifugally Atomized Powders”, J. Nucl. Mater., 321, 210–220 (2003) (Crys. Structure, Morphology, Phase Relations, Thermodyn., Experimental, Interface Phenomena, Kinetics, 31) Keiser, D.D., Jr., Clark, C.R., Meyer, M.K., “Phase Development in Al-Rich U-Mo-Al Alloys”, Scr. Mater., 51, 893–898 (2004) (Crys Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, 11) Leenaers, A., Van den Berghe, S., Koonen, E., Jarousse, C., Huet, F., Trotabas, M., Boyard, M., Guillot, S., Sannen, L., Verwerft, M., “Post-irradiation Examination of Uranium-7 wt% Molybdenum Atomized Dispersion Fuel”, J. Nucl. Mater., 335 (1), 39–47 (2004) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, 18) Tougait, O., Noel, H., “Stoichiometry of UAl4”, Intermetallics, 12, 219–223 (2004) (Crys. Structure, Experimental, 24) Schuster, J.C., “Al-Mo (Aluminium-Molybdenum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 30.12123.1.20, (2005) (Crys. Structure, Phase Diagram, Assessment, 61) Kim, Y.S., Hofman, G.L., Ryu, H.J., Hayes, S.L., “Irradiation-Enhanced Interdiffusion in the Diffusion Zone of U-Mo Dispersion Fuel in Al”, J. Phase Equilib. Diffus., 27(6), 614–621 (2006) (Morphology, Thermodyn., Calculation, Experimental, Interface Phenomena, Kinetics, 16) Ryu, H.J., Kim, Y.S., Hofman, G.L., Park, J.M., Kim, C.K., “Heats of Formation of (U,Mo)Al3 and U(Al, Si)3”, J. Nucl. Mater., 358, 52–56 (2006) (Morphology, Thermodyn., Experimental, 10) Ryu, H.J., Park, J.M., Kim, C.K., Kim, Y.S., Hofman, G.L., “Diffusion Reaction Behaviors of U-Mo/Al Dispersion Fuel”, J. Phase Equilib. Diffus., 27(6), 651–658 (2006) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, Kinetics, 21) Wieschalla, N., Bergmaier, A., Boeni, P., Boening, K., Dollinger, G., Grossmann, R., Petry, W., Roehrmoser, A., Schneider, J., “Heavy Ion Irradiation of U-Mo/Al Dispersion Fuel”, J. Nucl. Mater., 357, 191–197 (2006) (Morphology, Phase Relations, Experimental, Interface Phenomena, 9) Noel, H., Tougait, O., Potel, M., “Crystal Structures and Phase Stoichiometry of Nuclear Materials in the U-Si-C and U-Mo-Al Systems”, Collected Abstracts of the X International Conference on Crystal Chemistry of Intermetallic Compounds, Lviv, Ukraine, 17–20 September, 2007, Ivan Franko National University of Lviv, 5 (2007) (Crys. Structure, Experimental, 4) Soba, A., Denis, A., “An Interdiffusional Model for Prediction of the Interaction Layer Growth in the System Uranium-Molybdenum/Aluminum”, J. Nucl. Mater., 360, 231–241 (2007) (Morphology, Calculation, Theory, Interface Phenomena, Kinetics, 22) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

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Aluminium – Niobium – Nickel Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Lesley Cornish, Damian M. Cupid, Joachim Gro¨bner, Annelies Malfliet

Introduction Knowledge of the phase equilibria and thermodynamic properties of the Al-Nb-Ni system is relevant for many applications, especially for nickel-based superalloys. Pertinent to this is the order-disorder transformation between the disordered (Ni) and the ordered L12-Ni3Al [2003Du]. The excellent high temperature strength and high corrosion resistance of the binary intermetallic compound Ni3Al can be improved by alloying Nb to this compound, and this element has a large ternary solubility [2003Cer1]. The first experimental work on the Al-Nb-Ni system was by [1962Min1, 1962Min2] and [1965Kor], who both constructed the quasibinary section between Ni3Al and NbNi3. [1966Mar] determined the isothermal sections of most of the phase diagram at 900 and 1140˚C. Other investigations are mostly concentrated on the Ni rich corner [1970Cis, 1970Duv, 1979Oma, 1983Och, 1969Gus, 1989Hon]. The phase relations and homogeneity range of Ni3Al are studied by [1980Nas, 1994Jia]. An evaluation of the ternary Al-Nb-Ni system is made by [1993Sau]. [1992Lee, 2003Du, 2006Rag] reviewed the literature published up to this date. A Calphad type calculation of the entire phase diagram is reported by [2003Du], which is in good agreement with most of the experimental data. Although this calculation did not agree with all the experimental data, it has the best overall fit for describing the whole system. In cases where the agreement was less good, this was deduced to be because the samples had been annealed for relatively short times, considering the high melting points of most of the phases. Typically, at least 1200 h would be a reasonable annealing time for high temperature phases in bulk alloys with high melting compounds. Additionally, in all the isothermal sections calculated by [2003Du], the NbNi3 phase had an improbable shape in that there was some extension into the ternary on the Ni rich side, but not on the Nb rich side, and this has been smoothed in the present evaluation. An overview of the investigations considering phase equilibria, solid phases and thermodynamics of the system is given in Table 1.

Binary Systems The phase diagram of the binary Nb-Ni system is adopted from the calculations of [2006Che], which is in good agreement with most available experimental data. It is an extension of [2003Du], but used more recent experimental data. This diagram is redrawn in Fig. 1, but to be consistent with the other binary and ternary data the melting point of Nb is lowered from 2477 to 2469˚C. The binary Al-Ni phase diagram is taken from the MSIT evaluation by


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[2004Sal]. For the Al-Nb phase diagram, [Mas2] was accepted, with additional information on the Al rich liquidus Al-Nb was taken from [1973Wil].

Solid Phases The binary and ternary phases of the Al-Nb-Ni system are listed in Table 2. [1966Ben, 1968Hun] reported that the NbNi phase takes up a considerable quantity of Al. The binary phases, NbNi and Ni3Al have considerable extension into the ternary system [1966Ben]. [1968Hun] confirmed the large solubility of Al in the NbNi phase to be about 35 at.% Al. [1969Gus] measured a solubility of about 10 at.% Nb in Ni3Al, which only varied slightly over the temperature range 800 - 1200˚C. [1980Nas] confirmed this solubility at 1200˚C. The variation of the lattice parameter of Ni3Al along the line Ni3Al - NbNi3 was quite well established by [1962Min1, 1963Arb, 1969Gus, 1970Cis, 1980Nas, 1984Och1]. [1980Nas] reported a solubility of at least 5.5 at.% Nb in NiAl in equilibrium with τ1 and NbNi3 at 1200˚C. There are at least three true ternary phases τ1, τ2 and τ3 (denoted as H, L and M in the original literature) which exist only in the ternary diagram. The first data on these phases arose from [1964Sch, 1965Mar, 1965Ram]. They basically agree on the existence of the τ1, NbNi2Al, a Heusler type phase and τ2, Nb(Al,Ni)2 with a MgZn2 type structure. [1965Ram] reported that a phase, isostructural with Ti2Ni, exists in equilibrium with the τ1 phase, but no other studies report this phase. The τ2 phase has extensive solubility of Ni and Al confirming the formula Nb(AlxNi1–x)2, where x varies from 0.19 to 0.83 [1966Ben]. This result agrees with the data of [1966Mar]. In addition, above 900˚C an orthorhombic phase τ3 structurally related to the NbNi phase forms [1966Ben, 1967Sho, 1968Hun]. According to the phase diagram of [1966Ben], the τ3 phase extends from about 3 to 30 at.% Al. However, [1968Hun] observed a smaller homogeneity range for the τ3 phase. In addition, he found an additional ternary phase with higher Al content and like τ3, with a structure similar to NbNi. The discrepancy between the two authors could result from the difference in annealing time as [1966Ben] annealed only 20 h, while [1968Hun] used 168 h. However, [1968Hun] used only a limited number of samples, and the phase relations in his isothermal section are unusual. As no other experimental investigations report on this phase region, this part of the phase diagram is questionable. [1966Mar] did not observe the ternary τ3 phase at 900˚C and only a small extension of NbNi into the ternary. Thermal conductivity contours were used to determine the site preference of Nb in Ni3Al [2001Ter], and ab-initio calculations were used to investigate the site preference of Nb substitution in NiAl [2000Boz, 2001Son, 2002Boz]. Calculations based on the d-orbital level of the transition metals [1985Mor] as well as ab-initio methods [1991Eno] were used to predict the phase boundaries between NiAl and Ni3Al with dissolved Nb. First principles calculations were also used to calculate the relative stability of L12-Cu3Au and D0a-βCu3Ti structures of Ni3Al and NbNi3, each with additional Nb and Al respectively [1996Rav]. [2001Sav] used electrical resistivity measurements to examine the long range ordering of Ni3Al with Nb additions as a function of temperature. Order-disorder transition temperatures and the kinetics of the transition were studied [2001Sav].

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Quasibinary Systems [1965Kor] established that a simple quasibinary eutectic exists between NiAl and τ2 with stoichiometry NbNiAl, with the eutectic temperature and composition being respectively 1437˚C and 16.5 at.% Nb (Fig. 2). However, it should be noted that the thermal analysis data reported in tabular form give a 40˚C temperature range for the eutectic which would be very close to the maximal experimental error at that time. Thus, it is not fully conclusive with these data that the section is a true quasibinary. The position of the eutectic point does not coincide with that of the calculation of [2003Du], and in Fig. 6, the surface of solidification for τ2 has been moved to accommodate the experimental data. In addition, the solubility of Nb in the NiAl phase is less than might be expected from the measurements for NiAl in equilibrium with the τ1 phase [1980Nas], although the two measurements were from slightly different NiAl compositions.

Invariant Equilibria A ternary eutectic reaction (L Ð (Ni) + NiAl3 + NbNi3) was identified by [1973Kra] and [1973Lem] at 1270˚C and 13.6 at.% Nb and 79.1 at.% Ni. From this and the work of [1962Min1, 1962Min2, 1965Kor, 1969Nov] invariant reactions involving the liquid were deduced by [1992Lee]. [1992Lee] also included a possible ternary transition type solid state reaction NiAl + NbNi3 Ð NiAl3 + τ1. However, the evidence for this is considered too speculative here, and the reaction has not been included in Fig. 3, which is from the calculation of [2003Du]. [2003Rio] reported an invariant reaction L Ð NbAl3 + Nb2Al + τ2 at 1553˚C by DTA, with the ternary eutectic composition very near to the binary eutectic L Ð NbAl3 + Nb2Al. This is 150˚C higher than calculated by [2003Du], and the micrographs of [2003Rio] convincingly show three eutectic phases. Therefore, the calculated temperature is taken to be less reliable. The compositions of the phases participating in the invariant reactions are listed in Tables 3 and 4 as digitized from Fig. 4.

Liquidus, Solidus and Solvus Surfaces The liquidus of Ni rich alloys was determined by [1969Nov, 1973Kra, 1973Lem]. Based on these studies [1969Nov, 1973Kra, 1973Lem], [2003Du] undertook thermodynamic calculations and constructed a liquidus projection, which is shown in Fig. 4 [2003Du]. [2001Miu] investigated liquidus and solidus lines in the Ni rich corner between 1325 and 1450˚C (Figs. 5 and 6), which are consistent with the calculated results of [2003Du]. Both Fig. 5 and Fig. 6 are not adjusted to the accepted binary diagrams. The solvus line of (Ni) was reported by [1989Hon] and [2001Miu]. The solvus line of (Ni) [2001Miu] is shown in Fig. 7, and is qualitatively in agreement with [1989Hon].


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Isothermal Sections There have been a number of detailed studies of solid state equilibria in the Ni-rich region of the phase diagram: at 1300˚C by [1994Jia], at 1200˚C by [1962Min1, 1962Min2, 1970Cis, 1970Duv, 1980Nas, 1975Bok, 1994Jia, 1997Uey], at 1100˚C [1994Jia], at 1080˚C by [1975Bok], at 1000˚C by [1970Cis, 1983Och] at 800˚C by [1970Cis, 1969Gus, 1975Bok, 1994Jia], at 750˚C by [1970Duv] and at 600˚C by [1979Oma]. [1962Min1, 1962Min2, 1970Cis, 1970Duv, 1983Och] used both X-ray diffraction and metallography, while [1969Gus] used X-ray diffraction. A study by [1980Nas] used a variety of techniques which included metallography, X-ray diffraction and electron microprobe analysis. [1966Mar] using X-ray diffraction and [1966Ben] using both X-ray diffraction and metallography determined isothermal sections of almost the complete phase diagram at 900 and 1140˚C, respectively. In all the calculated sections, the NbNi3 phase showed some extension on the Ni rich side but not on the Nb rich side, and all of these have been smoothed here. The isothermal section at 1300˚C is given in Fig. 8. It is taken from the calculation by [2003Du] and in good agreement with the tie lines between (Ni) and Ni3Al measured by [1994Jia], although the agreement between NiAl3 and NiAl is less good, but acceptable. At 1200˚C (Fig. 9), [1962Min1, 1962Min2, 1970Cis, 1970Duv, 1975Bok, 1980Nas, 1994Jia, 1997Uey] show good agreement with [2003Du], although the (Ni) boundaries of [1970Cis] were slightly reduced, but within experimental error, especially considering that the samples have been annealed for only 65 h. [1980Nas] agrees with [2003Du], except for the minor phase analyses of two alloys which were probably compromised by the major phase (since the alloy composition was nearer to that of the major phase). In [1966Ben], the isothermal section at 1140˚C shows a three-phase field between τ2, NbAl3 and NiAl. This section is characterized by extensive solubility of the Ni3Al and NbNi phases in the ternary and the existence of three ternary compounds τ1, τ2 and τ3. In the calculated section [2003Du] (Fig. 10), the data point of NbAl3 [1966Ben] falls in a different phase field. However, this could be explained by the retention of NbAl3 due to a short annealing time (20 h), and so does not compromise the isothermal section of [2003Du]. At 1080˚C, the calculated isothermal section of [2003Du] (Fig. 11) is in good agreement with the experimental isothermal section of [1994Jia]. For the 1027˚C isothermal section, as shown in Fig. 12, the experimental data of [1983Och] show very good agreement with [2003Du], although Ni3Al boundary has a slight discrepancy. However, the experimental results of [1989Hon] for the three-phase region (Ni) + NbNi3 + Ni3Al do not fall in the calculated three-phase region of [2003Du], this could be because the alloys were only annealed for 555 h. The isothermal section at 900˚C of [2003Du] is given in Fig. 13. Most of the data at 900˚C [1966Mar] agree with [2003Du], except for the extension for NbNi for which [1966Mar] shows a very limited extent. The greater extent of NbNi might not have been observed due to the lack of equilibrium after annealing the alloys for only the relatively short time of 700 h (Nb having a high melting point would be slow to diffuse), and the high temperature Nb2Al phase might not have disappeared in this time. The extent of NbNi would fall in the NbNi-Nb2Al two-phase field of [1966Mar], and the identified three-phase fields could be due to a third phase which is still disappearing on annealing. Only partial isothermal sections were calculated at 800˚C, and these are shown in Figs. 14a and 14b. In the Ni rich corner at 800˚C (Fig. 14a), the experimental results for the position of the two-phase (Ni) + Ni3Al region [1969Gus, 1970Cis, 1994Jia] are not in agreement with the DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009

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calculation of [2003Du], but they were not used for the optimization. However, it must be pointed out that these three sets of experimental data agree well with each other, suggesting that the (Ni) actually has a lower solubility at 800˚C than is accepted by [2003Du]. Additionally, [1969Gus] shows that the (Ni) + Ni3Al two-phase field penetrates further into the ternary than that calculated by [2003Du]. The isothermal section at 800˚C in Fig. 14a was redrawn ensuring the NbNi3 phase is consistent with the accepted binary diagram. The Al rich corner at 800˚C (Fig. 14b) was investigated experimentally by [1965Mar] and the results are in fairly good agreement with the calculated work of [2003Du], even though they were not included in the optimization [2003Du]. Figure 15 shows the isothermal section at 750˚C. The experimental results of [1970Duv] in the Ni rich corner are in good agreement with each other, although (Ni) shows less solubility of Al and more for Nb than calculated by [2003Du]. Similarly, the isothermal section at 600˚C (given in Fig. 16) has some discrepancy between the experimental [1979Oma] and calculated results [2003Du], with the position of the two-phase Ni2Al3+NbAl3 region being slightly different. In general, the studies of [1962Min1, 1962Min2, 1969Gus, 1970Duv, 1980Nas, 1983Och] agree to within ±2 at.% Al and 1 at.% Nb on the composition of the three-phase equilibrium (Ni), Ni3Al and NbNi3, except [1980Nas] who reported up to 2 at.% Al solubility in NbNi3, while the other studies suggest that the solubility is very small. The results of [1970Cis] are quite different from these studies and have not been accepted here.

Temperature – Composition Sections A temperature-composition section at a constant 2.5 mass% Al showing the (Ni), Ni3Al and NbNi3 phases was produced by [1975Bok] experimentally, but is not consistent with the ternary diagrams accepted here (that of [2003Du] with some modifications in the liquidus surface) because of the shape of the liquidus near (Ni). [1962Min1, 1962Min2, 1965Kor] examined liquid/solid as well as solid state equilibria in two separate isopleths. [1962Min1, 1962Min2] examined alloys formed between Ni3Al and NbNi3 using thermal analysis, metallography, X-ray diffraction, hardness and resistivity measurement and established that a eutectic reaction exists between Ni3Al and NbNi3 at 1280˚C and 16 at.% Nb. This temperature-composition section is not a completely true quasibinary, since Ni3Al does not melt congruently. The temperatures of [1962Min1, 1962Min2] are close, although not in exact agreement, to the accepted Al-Ni binary, but are acceptable within experimental limits of that time. The composition of Ni3Al at the eutectic temperature was also determined by [1963Arb], and later [1969Tho], using metallography, reported that the eutectic composition is slightly lower in Nb at 15.4 at.% Nb. [1963Arb] carefully determined the solubility of Nb in Ni3Al at the quasibinary eutectic temperature using diffraction methods. Both of these additional studies are taken into account in Fig. 17 which has also been slightly altered from [1962Min1, 1962Min2] to be consistent with the AlNi binary system accepted here [2004Sal]. The calculated temperature-composition Ni3Al NbNi3 by [2003Du] is in good agreement with the experimental data of [1962Min1, 1963Arb, 1969Tho], except that the experimental results show less temperature dependence for Ni3Al.


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Thermodynamics A detailed thermodynamic modelling after the Calphad method was published by [2003Du]. Most of the available experimental information for the Al-Nb-Ni system was used for this calculation, except for those alluded to above [1965Mar, 1969Gus, 1970Cis, 1994Jia]. No fundamental discrepancies between the different experimental information taken from literature were observed. Experimental thermodynamic data were not available for this work. More recently [2006Hu] reported enthalpies of formation for two ternary phases τ1 and τ2 as well as other phases, and selected values are given in Table 5.

Notes on Materials Properties and Applications The most important application of the Al-Nb-Ni system is in the Ni-based superalloys, which can have up to 14 components. The matrix of this alloys comprised (Ni) and the major precipitates are based on Ni3Al. However, with increasing addition, the precipitates can be more complex. Nb is a useful addition for increasing the melting points. The Al-Nb-Ni system is also of interest as a metallic glass, with Al particles embedded in the glass phase, for structural materials. The glass formation and crystallization temperatures of amorphous (Ni60Nb40)100–xAlx alloys were investigated using differential scanning calorimetry [1987Akh, 2007Yu] and differential scanning calorimetry with differential thermal analysis [2004Lee]. [2007Yu] found that although the amorphous phase was resistant to attack by Al, the crystallized phase was not. Al87Ni10Nb3 was examined by [2004Aud], and He2+ ion bombardment [1991Ska], along with rapid quenching, and mechanical milling [2004Dia] were used to generate amorphous metal phases from NbNiAl and Nb45Ni50Al5 alloys. The Young’s modulus, yield strength, and ultimate tensile strength of an Al-particle reinforced Ni70Nb30 metallic glass was studied by [2006Yu]. [1998Far] produced NiAl matrix composites reinforced with Nb particles by reactive hot pressing of the elemental powders. Microhardness and compressive yield strength of the alloys were measured. [1995Net] used combustion synthesis (of elemental powders) to prepare selected Al-Nb-Ni alloys. [2000DeL] found that excess Al not participating in the aluminothermic reduction reaction for the formation of a Ni-65Nb (mass%) alloy dissolved into the NbNi phase, which is in agreement with the solubility of Al in NbNi. The effect on creep strength of tungsten and molybdenum layers applied to NiAl + Nb powders prior to pressing was studied by [2006Bel]. The ternary eutectic reactions in the Al-Nb-Ni system were used to produce directionally solidified (DS) eutectic alloys with in-situ, aligned composite microstructures without grain boundaries as potential candidate materials for replacement of Ni-based superalloys. After confirming the NbAl3-NbNiAl-Nb2Al ternary eutectic [2000Rio], the influence of growth rate on microstructure [2002Rio, 2004Rio, 2005Cos] and mechanical properties [2002Rio] were examined. [1992Rev] directionally solidified a NiAl-NbNiAl alloy near the eutectic maximum in the ternary. [1975Kra] directionally solidified an alloy near to the (Ni)-Ni3Al-NbNi3 eutectic, and high-temperature strengths of similarly solidified alloys in the same region were investigated by [1984Tor]. Dislocation structures and mechanical properties of a directionally solidified eutectic alloy with (Ni) / Ni3Al matrix with strengthening NbNi3 phase were investigated by [1978Sve]. DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009

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Investigations have been made on NiAl for potential shape memory effect [2002Kim], although the beneficial transformation was blocked by the transformation to Ni5Al3. However, Nb additions were found to facilitate the desired transformation. [1991Gra] investigated the oxidation behavior of two- and three-phase alloys in the Al-Nb-Ni system with thermogravimetric analysis. Oxidation of τ2 (NbNiAl) displayed quasilinear oxidation kinetics, whereas oxidation of ternary alloys of τ2 (NbNiAl) with additions of NiAl and (NiAl + Nb3Al) displayed parabolic oxidation rates. The corrosion behavior of Al-Nb-Ni alloys in mixed gas atmospheres of H2/H2O/H2S were generally parabolic, although linear behavior was also observed [1993He]. [1980Bhe] investigated the effect of different Al-Nb-Ni coating compositions on the oxidation of a selected Al-Nb-Ni alloy. Three types of oxidation behaviors were observed that depended on coating microstructure: formation of protective Al2O3 coatings, formation of non-protective Nb-oxide scales, and formation of non-protective NiO scales.

Miscellaneous Ni76Al22Nb22 and Ni3Al0.75Nb0.25 alloys were shown to obey an empirical relationship between bulk modulus B, molar volume V, and electron density η determined by [2004Li] as B = η2V. [2003Cer1] studied the effect of Nb concentration on interdiffusion coefficients in Ni3Al alloys with Nb using diffusion couples. Similarly, [2003Cer2] measured the interdiffusion coefficients of Nb and Al in the temperature interval 1173–1533 K with a Ni3Al + Ni-0.15Al0.075Nb couple to ascertain the effect of site preference on interdiffusion coefficients. The elastic modulus of a single crystal of composition Ni-0.19Al-0.06Nb was investigated using the velocity of elastic waves of frequencies 5.5, 15, and 50 MHz [2006Rin]. The coefficient of linear expansion was examined in several alloys based on the Ni3Al-NbNi3 composition line [1967Arb]. The results showed an inverse relationship between amount of NbNi3 and the linear expansion coefficient. The low-temperature (3.2 K-10.3 K) specific heat of NbNi2Al was determined using adiabatic calorimetry [1999Roc] and used to calculate the electronic specific heat, Debye temperature, and Einstein temperature. The experimental results were compared to theoretical calculations using the linear “muffin-tin” orbital-tight-binding (LMTO-TB) method. Other experimental work includes the microalloying Nb to NiAl [1991Sas], the characterization of the unstable growth interface with Widmansta¨tten-like precipitates between ternary Ni3Al and ternary NiAl [2001Kai], and Ni6NbAl metastable phase formation [1984Lia] through melt spinning and splat quenching experiments. The reversible hydrogen absorption properties at room temperature of NbNi and Nb10Ni9Al3 were also explored [2003Jou].


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. Table 1 Investigations of the Al-Nb-Ni Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied

Reference


[1962Arb]

XRD

Ni3Al-NbNi3 alloys

[1962Min1, 1962Min2]

Microscopy, thermal analysis, hardness, XRD

Vertical section Ni3Al-Ni3Nb

[1963Arb]

XRD, thermal analysis

Ni3Al-NbNi3 alloys at 1300˚C

[1964Sch]

XRD

Three ternary phases NbNi2Al, NbNiAl, NbNiAl2

[1965Gol]

Microchemical analysis, XRD

Ni3Al-NbNi3

[1965Kor]

Microscopy, thermal analysis, hardness, XRD

Vertical section NiAl to 40 at.% Nb

[1965Mar]

XRD, microstructure investigations NbNi2Al, NbNiAl annealed at 800˚C

[1965Ram]

XRD

NbNi2Al at 800˚C

[1966Ben]

Metallography, XRD

NbNiAl at 1140˚C

[1966Mar]

XRD

Complete composition range at 800 and 900˚C

[1967Sho]

XRD

Nb48Ni39Al13

[1968Hun]

XRD

Nb2Al region at 1000˚C

[1969Gus]

XRD

800, 1080 and 1200˚C in Ni rich range

[1969Nov]

Thermal analysis

Liquidus surface of Ni-Ni3Al-NbNi3 system

[1969Tho]

Metallography, chemical analysis

Ni3Al-NbNi3 alloys

[1970Cis]

Metallography, XRD and EPMA

Nickel rich corner at 800, 1000 and 1200˚C

[1970Duv]

Spectroscopy, light metallography, Ni rich corner (> 75 at.% Ni) at 750 and XRD 1200˚C

[1971Min]

XRD, microstructural analysis

Ni-Ni3Al-NbNi3 from room temperature to liquidus temperature

[1972Duv]

Optical metallography, TEM

750˚C Ni rich corner

[1973Kra, 1975Kra]

Directional solidification study, metallography, DTA

Ni-NbNi3-Ni3Al subsystem

[1973Lem]

Metallography, thermal analysis, XRD

Liquidus, Ni rich corner

[1973Wil]

Electronic structure calculation

Ni3Al at 1000˚C

[1975Bok]

Electron microscopy, solidification Ni-Ni3Al-NbNi3 at 800, 1080 and 1200˚C study

[1975Kau]

Calphad calculation

Complete composition range

[1978Gul]

Thermal analysis

Ni rich corner

[1979Oma]

Metallography, XRD

Complete system

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[1980Nas]

XRD, EPMA

Ni rich corner at 1200˚C

[1983Och]

Metallography, XRD

Solubility in Ni3Al at 1000˚C

[1984Arg]

XRD

Selected phases

[1984Bre]

Metallography, optical microscopy, Ni-Ni3Al-Nb Ni3 at 800, 1000 and 1200˚C SEM

[1984Och1]

XRD

(Ni), Ni3Al lattice parameters

[1984Och2]

Miedema calculation

Bond energies for Ni3Al with Nb

[1985Mis]

XRD

Ni3Al

[1985Tro]

XRD

τ2, NbNiAl (Laves phase)

[1986Bla]

XRD, metallography

NbNi2 from 800 to 1200˚C

[1988Mor]

Solidification study, XRD

Ni3Al at 1000˚C

[1989Hon]

DTA, optical microscopy, SEM

Solvus of Ni between 827 and 1227˚C

[1989Sub]

DTA, SEM, EPMA, XRD

Nb Al3 at 1300˚C

[1991Eno]

Cluster variation calculation

Ni and Ni3Al compositions below 1000˚C

[1991Ino]

Band structure calculations

NbAl3 with Ni substitutions

[1991Mis]

Metallography, XRD

(Ni) solvus surface

[1992Rev]

Metallography, SEM, EPMA

Al-16.5Nb-41.75Ni microstructures

[1994Jia]

EPMA, SEM

Ni rich corner at 800 and 1300˚C

[1997Uey]

SEM, EPMA, XRD

NbNi3 at 1200˚C

[2001Miu]

DTA

Ni rich corner from 900 to 1450˚C

[2002Kri]

XRD, DSC

Al-10Nb-50Ni and Al-10Nb-62Ni from room temperature to 700˚C

[2003Du]

Calphad calculation

Complete composition range above 600˚C

[2003Rio]

SEM, WDS, STA

Alloys: Al: 57.3–54.4; Nb: 41.8–33.3; Ni:0.9–12.3 (at.%)


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Comments/References

(Al) < 660.452

cF4 Fm 3m Cu

a = 404.96

pure Al at 25˚C [V-C2]

(Nb) < 2469

cI2 Im 3m W

a = 330.04

pure Nb at 25˚C [Mas2]

(Ni) < 1455

cF4 Fm 3m Cu

a = 352.40

pure Ni at 25˚C [Mas2]

NbAl3 < 1680

tI8 I4/mnm TiAl3

a = 384.1 c = 860.9

[1980Jor]

Nb2Al < 1940

tP30 P42/mnm σCrFe

a = 994.4 c = 517.2

66.7 at.% Nb [1980Jor]

Nb3Al < 2060

cP8 Pm 3n Cr3Si

a = 518.0

75 at.% Nb [1980Jor]

NiAl3 < 856

oP16 Pnma Fe3C

a = 661.3 b = 736.7 c = 481.1

[2004Sal]

Ni2Al3 < 1138

hP5 P3m1 Ni2Al3

a = 402.8 c = 489.1

36.8 to 40.5 at.% Ni [2004Sal]

β, NiAl < 1651

cP2 Pm 3m CsCl

a = 287

42 to 69.2 at.% Ni [2004Sal]

Ni5Al3 < 723

oC16 Cmmm Pt5Ga3

a = 753 b = 661 c = 376

63 to 68 at.% Ni [2004Sal]

γ´, Ni3Al < 1372

cP4 Pm 3m AuCu3

a = 356.77 a = 358.9 a = 356.32

73 to 76 at.% Ni [2004Sal]

NbNi < 1290

hR39 R3m W6Fe7

a = 489.4 c = 2674.0 a = 499.3 c = 2710.0

50 at.% Ni, 0 at.% Al [1968Hun] 24 at.% Ni, 30 at.% Al [1968Hun]

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. Table 2 (continued) Pearson Symbol/ Space Group/ Prototype



Comments/References

δ, NbNi3 < 1399

oP8 Pmmn TiCu3

a = 509.65 b = 423.23 c = 453.92

[1997Uey]

NbNi8

-

-

-

* τ1, NbNi2Al

cF16 Fm 3m BiF3

a = 597.0

25 at.% Al, 25 at.% Nb [1966Ben] ("H" phase)

* τ2, NbNiAl

hP12 P63/mmc MgZn2

a = 487.0 to 505.0 for formula Nb(AlxNi1–x)2 where x = 0.19 to c = 793.0 to 837.0 0.83 [1966Ben] ("L" phase)

* τ3, Nb10Ni9Al3

oP52 Pnma Nb10Ni9Al3

a = 930.3 b = 1626.6 c = 493.3

13 at.% Al, 48 at.% Nb [1967Sho], ("M" phase)

Metastable NbNi8 < 515

-

-

-

. Table 3 Invariant Four-Phase Equilibria Composition (at.%) Reaction

T [˚C]

Type

Phase

Al

L + Nb3Al Ð NbNi + (Nb)

1634

U1

L

Nb

Ni

29.00

51.00

20.00

L + Nb3Al Ð Nb2Al + NbNi

1633

U2

L

31.06

49.97

18.97

L+NbNi Ð NbNiAl + Nb10Ni9Al3

1547

U3

L

27.82

42.09

30.09

L + NbNi Ð Nb2Al + NbNiAl

1488

U4

L

49.18

38.11

12.71

L Ð NbAl3 +Nb2Al +NbNiAl

1408

E1

L

55.66

34.87

9.47

L Ð NbNi3 + Ni3Al + (Ni)

1267

E2

L

7.86

14.07

78.07

L + NiAl Ð NbNiAl + Ni3Al

1206

U5

L

16.28

21.56

62.16

L + NbNi3 Ð NbNiAl + Ni3Al

1205

U5

L

15.94

21.63

62.43

L + NbNiAl Ð Nb10Ni9Al3 + NbNi3

1203

U7

L

1.04

38.88

60.08

Ni3Al + NbNiAl Ð NiAl + NbNi3

1200

U8

-

-

-

-

L + Nb10Ni9Al3 Ð NbNi + NbNi3

1191

U9

L

0.58

40.21

59.21

NbNiAl + NiAl Ð NbNi2Al + NbNi3

1174

U10

-

-

-

-


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T [˚C]

Type

Phase

NiAl + NbNi3 ÐNbNi2Al + Ni3Al

1117

U11

-

L + NbNiAl Ð NiAl + NbAl3

1082

U12

L + NiAl Ð NbAl3 + Ni2Al3

1051

Nb10Ni9Al3 + NbNi3 Ð NbNiAl + NbNi Nb2Al + NbNi Ð NbNiAl + Nb3Al

Nb

Ni

-

-

-

L

63.62

11.59

24.79

U13

L

69.86

6.47

23.67

1019

U14

-

-

-

-

947

U15

-

-

-

-

L

-

-

-

-

-

-

L + Ni2Al3 Ð NbAl3 + NiAl3

849

U16

Nb3Al + NbNi Ð NbNiAl + (Nb)

797

U17

NbNi + (Nb) Ð NbNi2Al + NbNiAl

779

U18

NbNiAl + (Nb) Ð NbNi2Al + Nb3Al

728

U19

L + NbAl3 Ð NiAl3 + (Al)

647

U20

Al

-

L

. Table 4 Invariant Three-Phase Equilibria Composition (at.%) Reaction

T [˚C]

Type

Phase

Al

L + Nb3Al Ð NbNi

1635

p3

L

L + NbNi Ð Nb10Ni9Al3

1615

p4

L Ð NbNi + NbNiAl

1571

L Ð NiAl + NbNiAl

Nb

Ni

29.68

50.56

19.76

L

24.6

45.10

30.30

e1

L

37.42

39.09

23.49

1460

e3

L

37.06

18.87

44.07

L Ð NbAl3 + NbNiAl

1419

e4

L

57.5

31.25

11.25

L Ð NbNi3 + NbNiAl

1276

e7

L

7.28

29.96

62.76

L Ð NbNi3 + Ni3Al

1275

e8

L

11.32

15.44

73.24

L Ð NbNiAl + Ni3Al

1204

e9

-

-

-

-

NiAl + NbNiAl Ð NbNi2Al

1178

p7

-

-

-

-

Nb10Ni9Al3 Ð NbNiAl +NbNi

910

e11

-

-

-

-

NbNi Ð NbNi2Al + (Nb)

815

e12

-

-

-

-

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. Table 5 Thermodynamic Properties of Single Phases Temperature Range [˚C]

Phase

Property, per mole of atoms [J, mol, K]

Comments

τ1, Al25Ni50Nb25

25

–38600 ± 1100

Enthalpy of formation [2006Hu]

τ2, Al27Ni40Nb33:

25

–39600 ± 1600


τ2, Al33.3Ni33.3Nb33.3

25

–40900 ± 1500


τ2, Al40Ni20Nb33: (N.B. not 100%)

25

–46200 ± 1300


NiAl3 of composition: Al75Ni17Nb8

25

–40300 ± 1100


Ni3Al of composition: Al20Ni76Nb4

25

–28400 ± 800


. Table 6 Investigations of the Al-Nb-Ni Materials Properties Reference


Type of Property

[1967Arb]

Hardness, mechanical testing

Vickers Hardness, Young’s modulus, thermal expansion

[1978Sve, 2006Gre]

TEM

Dislocation structures

[1980Bhe, 1991Gra]

Thermogravimetric analysis

Oxidation behavior, oxidation rates

[1984Tor]

Metallography, directional solidification, mechanical testing

Phases, high temperature strength, fracture toughness

[1991Laa]

HIP poweder compacts, mechanical testing

Young’s modulus, Poisson’s ratio, and fracture toughness at RT

[1991Och]

Compression testing

Stress-strain behavior

[1992Gra]


Oxidation resistance of intermetallic phases

[1992Rev]

Directional solidification, microhardness

Microhardness and fracture strength

[1993He]


Corrosion behavior

[1993Rei]

Compression testing

High temperature strength, fracture toughness

[1996Mac]

Mechanical testing

Deformation behaviour of τ2, NbNiAl


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[1998Far]

Solid state sintering of elemental powders; Hardness testing, compression testing

NiAl-Nb composites; Hardness testing, compression testing

[1999Roc]

Calorimetry

Low temperature specific heat

[2002Kim]

Metallography, DTA

Phase transformations

[2002Rio]

Hardness testing

Hardness, fracture toughness

[2003Cer1, 2003Cer2]

Diffusion couples

Interdiffusion coefficients

[2006Bel]

Creep testing

Creep strength

[2006Hag]

Three-point bending

Fracture toughness

[2006Rin]

Velocity of elastic waves

Elastic modulus

[2006Yu]

Compression testing

Fabrication and characterization of Ni-Nb particle-reinforced Al-based composite

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Al–Nb–Ni

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. Fig. 1 Al-Nb-Ni. Adopted binary Nb-Ni system


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Al–Nb–Ni

. Fig. 2 Al-Nb-Ni. Experimental quasibinary system between NiAl and τ2 ( NbNiAl)

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. Fig. 3a Al-Nb-Ni. Reaction scheme, part 1

Al–Nb–Ni


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Al–Nb–Ni

. Fig. 3b Al-Nb-Ni. Reaction scheme, part 2

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. Fig. 3c Al-Nb-Ni. Reaction scheme, part 3

Al–Nb–Ni


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. Fig. 4 Al-Nb-Ni. Liquidus surface projection

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. Fig. 5 Al-Nb-Ni. Experimental liquidus lines in the Ni-rich corner between 1325 and 1450˚C


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Al–Nb–Ni

. Fig. 6 Al-Nb-Ni. Experimental solidus lines in the Ni-rich corner between 1325 and 1450˚C

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. Fig. 7 Al-Nb-Ni. Dependence of the solvus line of (Ni) on Al concentration


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Al–Nb–Ni

. Fig. 8 Al-Nb-Ni. Isothermal section at 1300˚C

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Al–Nb–Ni

. Fig. 14a Al-Nb-Ni. Isothermal section at 800˚C, Ni rich corner

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. Fig. 14b Al-Nb-Ni. Isothermal section at 800˚C, Al rich corner


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. Fig. 17 Al-Nb-Ni. Experimental temperature - composition section Ni3Al - NbNi3

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References [1962Arb]

[1962Min1]

[1962Min2]

[1963Arb] [1964Sch] [1965Gol]

[1965Kor]

[1965Mar]

[1965Ram]

[1966Ben] [1966Mar]

[1967Arb]

[1967Sho] [1968Hun]

[1969Gus]

[1969Nov]

[1969Tho]

[1970Cis] [1970Duv]

Arbuzov, M.P., Chuprina, V.G., “X-Ray Determination of the Crystal Structures of Alloys in the Ni3AlNi3Nb System” (in Russian), Akad. Nauk SSSR, Inst. Metall. Im A.A. Baikova, (8), 85–87 (1962) (Crys. Structure, Experimental, 13) Mints, R.S., Belyaeva, G.F., Malkov, Yu.S., “The Reaction Between the Metallic Compounds Ni3Al and Ni3Nb” (in Russian), Dokl. Akad. Nauk SSSR, 143(4), 871–874 (1962) (Phase Relations, Phase Diagram, *, 26) Mints, R.S., Belyaeva, G.F., Malkov, Y.S., “Equlibrium Diagram of the Ni3Al-Ni3Nb System”, Russ. J. Inorg. Chem. (Engl. Transl.), 7, 1236–1239 (1962), translated from Zh. Neorg. Khim., 7(10), 2382–2385 (1962) (Experimental, Morphology, Phase Relations, *, 32) Arbuzov, M.P., Chuprina, V.G., “Ageing Alloys of the System Ni3Al-Ni3Nb” (in Russian), Izv. V. U. Z., Fiz., 5, 82–85 (1963) (Crys. Structure, Phase Diagram, Experimental, *, 2) Schubert, K., Meissner, H.G., Raman, A., Rossteutscher, W., “Structural Data of Some Metallic Phases”, Naturwissenschaften, 51, 287 (1964) (Crystal Structure, 0) Golubtsova, R.B., “Selective Isolation of Metallic Compounds from an Alloy of the Ni3Al-Ni3Nb System” (in Russian), Dokl. Akad. Nauk SSSR, 160, 1311–1314 (1965) (Experimental, Phys. Prop., Kinetics, 13) Kornilov, I.I., Mints R.S., Guseva, L.N., Malkov, Y.S., “Interaction Between the Compound NiAl and Niobium”, Russ. Metall. (Engl. Transl.), (6), 93–96 (1965), translated from Izv. Akad. Nauk SSSR, Met., (6), 132–136 (1965) (Experimental, Morphology, Phase Diagram, Phase Relations, #, 5) Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of the MnCu2Al and MgZn2 Types Containing Aluminum and Gallium”, Sov. Phys. Crystallogr., 9, 619–620 (1965) translated from Kristallografiya, 9, 737–738 (1964) (Crys. Structure, Experimental, 4) Raman, A., Schubert, K., “On the Crystal Structure of Some Alloy Phases Related to TiAl3. III. Investigations in Several T-Ni-Al and T-Cu-Al Systems” (in German), Z. Metallkd., 56, 99–104 (1965) (Crys. Structure, Experimental, Phase Diagram, 14) Benjamin, J.S., Giessen, B.C., Graut, N.J., “Intermediate Phases in the Ternary System Nb(Cb)-Ni-Al at 1140˚C”, Trans. Metall. Soc. AIME, 236, 224–226 (1966) (Experimental, Phase Relations, #, 8) Markiv, V.Ya., Matushevskaya, N.F., Kuz’ma, Yu.B., “X-Ray Diffraction of the Nb-Ni-Al System”, Russ. Metall. (Engl. Transl.), (6), 72–74 (1966), translated from Izv. Akad. Nauk SSSR, Met., (6), 127–129 (1966) (Experimental, Phase Relations, *, 2) Arbuzov, M.P., Chuprina, V.G., “Determination of Some Physical Properties of the Ni3Al-Ni3Nb Alloys”, Russ. Metall. (Engl. Transl.), (3), 78–80 (1967), translated from Izv. Akad. Nauk SSSR, Met., (3), 174–179 (1967) (Experimental, Phase Relations, Phys. Prop., 7) Shoemaker, C.B., Shoemaker, D.P., “The Crystal Structure of the M Phase, NbNiAl”, Acta Crystallogr., 23, 231–238 (1967) (Crys. Structure, Experimental, 18) Hunt, C.R. Jr., Raman, A., “Alloy Chemistry of σ (βαffl U)-Related Phases. I. Extens Ion of μ- and Occurrence of μ’-Phases in the Ternary Systems Niobium(Tantalum)-X-Aluminum X = Iron, Cobalt, Nickel, Copper, Chromium, Molybdenum)”, Z. Metallkd., 59(9), 701–707 (1968) (Crys. Structure, Phase Diagram, *, 14) Guseva, L.N., Mints, R.S., Malhov, Y.S., “Phase Equilibria in the Ni-Ni3Al-Ni3Nb System at 800–1200˚C”, Russ. Metall. (Engl. Transl.), (5), 120–122 (1969), translated from Izv. Akad. Nauk SSSR, Met., (5), 186–188 (1969) (Experimental, Phase Relations, *, 12) Novic, F.S., Mints, R.S., Malkov, Yu.S., “Plotting the Liquidus Surface of a Ni-Ni3Al-Ni3Nb Ternary System” (in Russian) in “Teor. Eksp. Metody Issled. Diagramm Sostoyaniya Metal. Sistem”, Ageev, N.V. (Ed.), Nauka, Moscow, 145–150 (1969) (Phase Diagram, Experimental, *, 10) Thompson, E.R., Lemkey, F.D., “Structure and Properties of Ni3Al Eutectic Alloys Produced by Unidirectional Solidification”, Trans. Quart., A.S.M., 62, 140–154 (1969) (Phase Diagram, Experimental, *, 35) Cisse, J., Davies, R.G., “Nickel-Rich Portion of the Nickel-Aluminum-Niobium Phase Diagram”, Met. Trans., 1, 2003–2006 (1970) (Experimental, Mechan. Prop., Morphology, Phase Diagram, 9) Duvall, D.S., Donachie, M.J., “Phase Equilibria in Nickel-Rich Ni-Al-Nb Alloys”, J. Inst. Met., 96, 182–187 (1970) (Experimental, Morphology, Phase Diagram, *, 16)


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11 [1971Min]

[1972Duv] [1973Kra]

[1973Lem]

[1973Wil] [1975Bok]

[1975Kau] [1975Kra]

[1978Gul] [1978Sve]

[1979Oma]

[1980Bhe]

[1980Jor]

[1980Nas] [1983Och] [1984Arg] [1984Bre] [1984Lia] [1984Och1]

[1984Och2] [1984Tor]

Al–Nb–Ni Mints, R.S., Malkov, Y.S., Davydor, N.I., Kofanova, G.N., “Polythermal Sections of the Ni-Ni3Al-Ni3Nb System at Nb 5 and 10 wt.-%” (in Russian), Diagrammy Sost. Metall. Sistem, 158–160 (1971) (Phase Diagram, Phase Relations, Experimental, Mechan. Prop., 8) Duvall, D.S., Donachie, M.J. Jr., “Precipitation Characteristics of Nickel-Rich Ni-Al-Nb Alloys”, J. Inst. Met., 100, 6–12 (1972) (Experimental, Morphology, Phase Relations, 22) Kraft, E.H., Thompson, E.R., “Directional Solidification of Ni-Nb, Ni-Al-Nb, Ni-Cr-Nb and Ni-Cr-AlNb Alloys”, Proc. Conf. “In Situ Composites”, National Academy of Science, Washington, 297–303 (1973) (Phase Relations, Experimental, *, 18) Lemkey, F.D., Thompson, E.R., “Eutectic Superalloys Strengthened by Aligned S-Ni3Nb Lamellas, γ’(Ni3Al) Precipitates and Reduced Interlamellar Spacing”, Proc. Conf. “In Situ Composites”, National Academy of Science, Washington, 161 (1973) (Phase Relations, Experimental) Willey, L.A., “The Al-Nb System” in “Metals Handbook”, 8th Ed., Vol. 8, ASM, Metals Park Ohio, 342 (1973) (Phase Relations, Review, #, 3) Bokshteyn, S.Z., Vasilenok, L.B., Kishkin, S.T., Nazarova, M.P., Svetlov, I.L., Sorokina, L.P., Khusnetdinov, F.M., “Formation of the Eutectic γ/γ’ - δ Microstructure During Directional Crystallization and After Heat Treatment”, Phys. Met. Metallogr., 79–85 (1975) (Experimental, Morphology, Phase Diagram, 17) Kaufman, L., Nesor, H., “Calculation of Superalloy Phase Diagrams. Part III”, Metall. Trans. A, 6(11), 2115–2122 (1975) (Calculation, Crys. Structure, Phase Diagram, Thermodyn., 35) Kraft, E.H., Thompson, E.R., “Directional Solidification of Ni-Nb, Ni-Al-Nb, Ni-Cr-Nb and Ni-Cr-AlNb Alloys”, 2nd Conf. on in Situ Composites, Boston, 297–307 (1975) (Experimental, Thermodyn., Phase Relations, 18) Gulyaev, B.B., Grigorash, E.F., Efimova, M.N., “Solidification Range of Ni Alloys”, Met. Sci. Heat Treat., 20, 914–917 (1978) (Experimental, Phase Relations, 8) Svetlov, I.L., Vasilenok, L.B., Khusnetdinov, F.M., Sidorov, V.V., Nazarova, M.P., “Plastic Deformation of Directionally Crystallized Eutectic Ni/Ni3Al-Ni3Nb”, Phys. Met. Metallogr., 45(1), 124–129 (1978) (Experimental, Mechan. Prop., 12) Omarov, A.K., Smagulov, S.U., Askarov, M.S., “Study of Phase Equilibriums in Alloys of Al-Ni-Zn and Al-Nb-Ni Ternary Systems” (in Russian), Metall. Obogashch., 133–137 (1979) (Abstract, Experimental, Phase Diagram, Phase Relations, 5) Bhedwar, H.C., Heckel, R.W., Laughlin, D.E., “The Oxidation Behavior of Aluminide-Coated γ/δ Directional Eutectics”, Metall. Trans. A, 11, 1303–1314 (1980) (Experimental, Mechan. Prop., Morphology, Phase Diagram, Phase Relations, Phys. Prop., 25) Jorda, J.L., Fluekiger, R., Mueller, J., “A New Metallurgical Investigation of the NiobiumAluminium System”, J. Less-Common Met., 75, 227–239 (1980) (Crys. Structure, Phase Diagram, Experimental, #, 20) Nash, P., Kavishe, F.P.L., West, D.R.F., “Nickel-Rich Region of Ni-Al-Nb System at 1473 K”, Met. Sci., 14(4), 147–149 (1980) (Experimental, Phase Relations, Crys. Structure, *, 21) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Additions”, Bull. P.M.E. (T.I.T.), 52, 1–17 (1983) (Phase Diagram, Experimental, 45) Argent, B.B., “Phase Diagrams of Alloys Based on Niobium”, Metals Society/AIME Acc. No. 84(7), 72–486, 325–415 (1984) (Crys. Structure, Phase Diagram) Brezovsky, M., “Morphology of Eutectic Composite Material γ-γ´”, Kovove Mater., 1 (22), 104–112 (1984) (in Slovakija) (Experimental, 17) Liang, W.W., Standley, R., Nash, P., Skowron, M., “The Relative Stabilities of the Ni6AlX (X = V, Nb, Ta) Phases”, J. Mater. Sci. Lett., 3(3), 259–261 (1984) (Experimental, 5) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni (γ), Ni3Al (γ’) and Ni3Ga (γ’) Solid Solutions”, Bull. Res. Lab. Precis. Machin. Electron., Tokyo Inst. Technol., 53, 15–28 (1984) (Crys. Structure, Experimental, *, 66) Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32(2), 289–298 (1984) (Experimental, Phase Diagram, 90) Toropov, V.M., Bondarenko, Yu.A., “Properties of High-Temperature Alloys of the System Ni-Al-Nb with a Unidirectional Eutectic Structure”, Met. Sci. Heat Treat., 26(9-10), 660–664 (1984), translated from Metalloved. Term. Obrab. Met., (9), 11–15, 1984 (Experimental, Mechan. Prop., Phase Diagram, 8)

DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009

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Al–Nb–Ni [1985Mis]

[1985Mor] [1985Tro] [1986Bla] [1987Akh] [1988Mor]

[1989Hon] [1989Sub]

[1991Eno]

[1991Gra]

[1991Ino] [1991Laa]

[1991Mis] [1991Och]

[1991Sas]

[1991Ska]

[1992Gra]

[1992Lee] [1992Rev]

[1993He]

[1993Rei]

11

Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni(γ), Ni3Al(γ’) and Ni3Ga(γ’) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33(6), 1161–1169 (1985) (Experimental, 64) Morinaga, M., Yukawa, N., Ezaki, H., Adachi, H., “Solid Solubilities in Nickel-Based F.C.C. Alloys”, Philos. Mag. A, 51(2), 247–252 (1985) (Phase Diagram, Phase Relations, Experimental, 26) Trojko, R., Blazina, Z., “Metal-Metalloid Exchange in Some Friauf-Laves Phases Containing Two Transition Metals”, J. Less-Common Met., 106, 293–300 (1985) (Crys. Structure, Experimental, 13) Blazina, Z., Trojko, R., “Structural Investigations of the Nb(1–x)SixT2 and Nb(1–x)AlxT2 (T = Cr, Mn, Fe, Co, Ni) Systems”, J. Less-Common Met., 119, 297–305 (1986) (Crys. Structure, Experimental, 6) Akhtar, D., Misra, R.D.K., “Formation and Stability of Ni60Nb40-xAlx”, J. Mater. Sci. Lett., 6(1), 29–30 (1987) (Experimental, 19) Morinaga, M., Sone, K., Kamimura, T., Ohtaka, K., Yukawa, N., “X-Ray Determination of Static Displacements of Atoms in Alloyed Ni3Al”, J. Appl. Crystallogr., 21, 41–46 (1988) (Crys. Structure, Experimental, 18) Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of γSolvus in Ni-Al-X Ternary Systems”, Mater. Res. Soc. Symp. Proc., 133, 429–440 (1989) (Experimental, Phase Diagram, Phase Relations, 35) Subramanian, P.R., Simmons, J.P., Mendiratta, M.G., Dimiduk, D.M., “Effect of Solutes on Phase Stability in Al3Nb”, Mater. Res. Soc. Symp. Proc., 133(3), 51–56 (1989) (Experimental, Mechan. Prop., Phase Diagram, 12) Enomoto, M., Harada, H., Yamazaki, M., “Calculation of γ’/γ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2), 143–158 (1991) (Assessment, Calculation, Phase Diagram, 34) Grabke, H.J., Steinhorst, M., Brumm, M., Wiemer, D., “Oxidation and Intergranular Disintegration of the Aluminides NiAl and NbAl3 and Phases in the System Nb-Ni-Al”, Oxid. Met., 35(3-4), 199–222 (1991) (Electronic Structure, Experimental, Kinetics, Thermodyn., 30) Inoue, H.R.P., Kitamura, M., Wayman, C.M., Chen, H., “Phase Stability of Al3Nb as a Function of Nickel Additions”, Philos. Mag. Lett., 63(6), 345–353 (1991) (Crys. Structure, Experimental, 19) Laag, R., Kaysser, W.A., Petzow, G., “A Comparative Study on the Influence of Nb and Ti Additions to Different Processed Atomized NiAl Powders”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 821–826 (1991) (Experimental, Mechan. Prop., 9) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the γ Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123–130 (1991) (Assessment, Experimental, Phase Diagram, 5) Ochiai, S., Shirokura, T., Doi, Y., Kojima, Y., “Microstrucutres and Mechanical Properties of Ni-Nb Aluminides Produced by MA Process”, ISIJ Int., 31(10), 1106–1112 (1991) (Experimental, Mechan. Prop., 27) Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on thr Solidified Structure of NiAl”, Proc. Conf. Intermetal. Comp. - Struct. Mechan. Prop., 877–881 (1991) (Abstract, Experimental, Mechan. Prop., Phase Diagram, 10) Skakov, Yu.A., Djakonova, N.P., Edneral, N.V., Koknaeva, M.R., Semina, V.K., “Some Peculiarities of the Atomic Structure of Metallic Phase Formed During Liguid Quenching and Solid State Reactions”, Mater. Sci. Eng. A, 133, 560–564 (1991) (Experimental, Phase Relations, 5) Grabke, H.J., Brumm, M., Steinhorst, M., “Development of Oxidation Resistant High Temperature Intermetallics”, Mater. Sci. Technol., 8, 339–344 (1992) (Experimental, Interface Phenomena, Phase Diagram, Thermodyn., 21) Lee, K.J., Nash, P., “The Al-Nb-Ni System”, to be published in J. Phase Equilib., (Crys. Structure, Phase Diagram, Review, #, 27) a copy is available at MSI Reviere, R.D., Noebe, R.D., Oliver, B.F., “Processing, Microstructure and Low-Temperature Properties of Directionally Solidified NiAl/NiAlNb Alloys”, Mater. Lett., 14(2-3), 149–155 (1992) (Morphology, Phase Relations, Mechan. Prop., 19) He, Y.R., Douglass, D.L., “The Corrosion Behavior of Ni-Al Alloys and Ni-Nb-Al Alloys in a H2/H2O/ H2S Gas Mixture”, Oxid. Met., 40(3-4), 337–371 (1993) (Experimental, Kinetics, Morphology, Phase Relations, 28) Reip, C.-P., Sauthoff, G., “Deformation Behaviour of the Intermetallic Phase Al3Nb with DO22 Structure and of Al3Nb-base Alloys: Part 1. Physical Properties and Short-Term Behaviour”, Intermetallics, 1, 159–169 (1993) (Experimental, Phase Diagram, Phys. Prop., 30)


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38

11 [1993Sau]

[1994Jia]

[1995Net] [1996Mac] [1996Rav]

[1997Uey]

[1998Far]

[1999Roc]

[2000Boz] [2000DeL]

[2000Rio] [2001Kai]

[2001Miu]

[2001Sav]

[2001Son] [2001Ter]

[2002Boz] [2002Kim]

[2002Kri]

[2002Rio]

Al–Nb–Ni Saunders, N., “Aluminium - Niobium - Nickel”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.10206.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, XX) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between γ (A1), γ’ (L12) and β (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473–485 (1994) (Crys. Structure, Experimental, Phase Diagram, 25) Neto, R.M.L., Ferreira, P.I., “Reaction Sintering of Nb-Ni-Al Intermetallic Alloys”, Mater. Sci. Eng. A, 193, 549–555 (1995) (Experimental, Kinetics, Mechan. Prop., Phase Relations, Thermodyn., 18) Machon, L., Sauthoff, G., “Deformation Behaviour of Al-Containing C14 Laves Phase Alloys”, Intermetallics, 4, 469–481 (1996) (Experimental, Phase Relations, 41) Ravindran, P., Subramoniam, G., Asokamani, R., “Ground-State Properties and Relative Stability Between the L12 and Doa Phases of Ni3Al by Nb Substitution”, Phys. Rev. B, 53(3), 1129–1137 (1996) (Calculation, Crys. Structure, Thermodyn., 44) Ueyama, T., Ghanem, M.M., Miura, N., Takeyama, M., Matsuo, T., “Phase Stability of Ni3Nb-δ Phase in Ni-Nb-M Systems at Elevated Temperatures”, THERMEC´97, Intern. Conf. Thermomechan. Proc. Steels Other Mater., TMS, Warrendale, USA, 2, 1753–1760 (1997) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 16) Farber, L., Gotman, I., Gutmanas, E.Y., Lawley, A., “Solid State Synthesis of NiAl-Nb Composites from Fine Elemental Powders”, Mater. Sci. Eng. A, 244(1), 97–102 (1998) (Experimental, Morphology, Phase Relations, 29) da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., da Silva, C.M., Gomes, A.A., “Specific Heat and Electronic Structure of Heusler Compounds Ni2Tal (T = Ti, Zr, Hf, V, Nb, Ta)”, Physica B (Amsterdam), 269, 154–162 (1999) (Crys. Structure, Experimental, 20) Bozzolo, G., Noebe, R.D., Honecy, F., “Modeling of Ternary Element Site Substitution in NiAl”, Intermetallics, 8, 7–18 (2000) (Crys. Structure, Review, 34) De Lima, B.B., Ramos, A.S., Nunes, C.A., Conte, R.A., “Ni-65 wt.% Nb Alloy by Aluminothermic Reduction Process”, Int. J. Refract. Met. Hard Mater., 18, 267–271 (2000) (Experimental, Phase Diagram, 6) Rios, C.T., Milenkovic, S., Caram, R., “Directional Growth of Al-Nb-X Eutectic Alloys”, J. Cryst. Growth, 211, 466–470 (2000) (Experimental, Phase Relations, 8) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of γ´/β Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168–175 (2001) (Experimental, Phase Relations, Thermodyn., 21) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: V, Nb and Ta) Ternary Systems”, J. Phase Equilib., 22, 345–351 (2001) (Experimental, Phase Relations, 9) Savin, O.V., Stepanova, N.N., Akshentsev, Yu.N., Rodionov, D.P., “Ordering Kinetics in Ternary Ni3Al-X Alloys”, Scr. Mater., 45(8), 883–888 (2001) (Crys. Structure, Electr. Prop., Experimental, Kinetics, Thermodyn., 18) Song, Y., Guo, Z.X., Yang, R., Li, D., “First Principles Study of Site Substitution of Ternary Elements in NiAl”, Acta Mater., 49, 1647–1654 (2001) (Calculation, Electronic Structure, 17) Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurement”, J. Mater. Res., 16(8), 2314–2320 (2001) (Calculation, Crys. Structure, Experimental, Transport Phenomena, 63) Bozzolo, G.H., Noebe, R.D., Amador, C., “Site Occupancy of Ternary Additions to B2 Alloys”, Intermetallics, 10, 149–159 (2002) (Crys. Structure, Review, 27) Kim, S.-H., Oh, M.-H., Wee, D.-M., “Phase Transformation Behavior and Critical Temperature for Operation of Ni-Al Shape Memory Alloys Including Ternary Elements” (in Korean), J. Korean Inst. Met., 40(6), 621–627 (2002) (Experimental, Morphology, Phase Relations, 18) Krivoroutchko, K.A., Kulik, T., Fadeeva, V.I., Portnoy V.K., “Formation of Stable and Metastable Phases in Ni-Al-Nb and Ni-Al-Me-C (Me = Ti, Nb or V) Powder Systems during Mechanical Alloying and Thermal Treatment”, J. Alloys Compd., 333, 225–230 (2002) (Crys. Structure, Experimental, 13) Rios, C.T., Milenkovic, S., Gama, S., Caram, R., “Influence of the Growth Rate on the Microstructure of a Nb-Al-Ni Ternary Eutectic”, J. Cryst. Growth, 237-239, 90–94 (2002) (Experimental, Phase Relations, 8)

DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009

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Al–Nb–Ni [2003Cer1]

[2003Cer2]

[2003Du]

[2003Jou]

[2003Rio] [2004Aud]

[2004Dia]

[2004Lee]

[2004Li]

[2004Rio]

[2004Sal]

[2005Cos]

[2006Bel]

[2006Che] [2006Gre]

[2006Hag]

[2006Hu] [2006Rag]

11

Cermak, J., Rothova, V., “Concentration Dependence of Ternary Interdiffusion Coefficirnts in Ni3Al/ Ni3Al-X Couples with X = Cr, Fe, Nb and Ti”, Acta Mater., 51(15), 4411–4421 (2003) (Electronic Structure, Experimental, Transport Phenomena, 15) Cermak, J., Gazda, A., Rothova, V., “Interdiffusion in Ternary Ni3Al/Ni3Al-X Diffusion Couples with X = Cr, Fe, Nb and Ti”, Intermetallics, 11(9), 939–946 (2003) (Experimental, Kinetics, Transport Phenomena, 24) Du, Y., Chang, Y.A., Gong, W., Huang, B., Xu, H., Jin, Zh., Zhang, F., Chen, S.-L., “Thermodynamic Properties of the Al-Nb-Ni System”, Intermetallics, 11(1-2), 995–1013 (2003) (Assessment, Phase Diagram, Thermodyn., 45) Joubert, J.-M., Pommier, C., Leroy, E., Percheron-Guegan, A., “Hydrogen Absorption Properties of Topologically Close-Packed Phases of the Nb-Ni-Al System”, J. Alloys Compd., 356-357, 442–446 (2003) (Crys. Structure, Experimental, 17) Rios, C.T., Milenkovic, S., Caram, R., “A Novel ternary Eutectic in the Nb-Al-Ni System”, Scr. Mater., 48(10), 1495–1500 (2003) (Crys. Structure, Experimental, Kinetics, Morphology, Phase Relations, 10) Audebert, F., Mendive, C., Vidal, A., “Structure and Mechanical Behaviour of Al-Fe-X and Al-Ni-X Rapidly Solidified Alloys”, Mater. Sci. Eng. A, 375-377, 1196–1200 (2004) (Electronic Structure, Experimental, 22) Diakonova, N.P., Sviridova, T.A., Semina, V.K., Skakov, Yu.A., “Intermetallic Phase Stability on High Energy Treatments (Rapid Quenching, Ion Irradiation and Mechanical Milling)”, J. Alloys Compd., 367, 199–204 (2004) (Crys. Structure, Experimental, Phase Relations, 18) Lee, M.H., Kim, W.T., Kim, D.H., Kim, Y.B., “The Effect of Al Addition on the Thermal Properties and Crystallization Behavior of Ni60Nb40 Metallic Glass”, Mater. Sci. Eng. A, 375-377, 336–340 (2004) (Crys. Structure, Experimental, Interface Phenomena, 12) Li, C., Chin, Y.L., Wu, P., “Correlation between Bulk Modulus of Ternary Intermetallic Compounds and Atomic Properties of their Constituent Elements”, Intermetallics, 12, 103–109 (2004) (Electronic Structure, Thermodyn., 24) Rios, C.T., Oliveira, M.F., Caram, R., Botta F.W.J., Bolfarini, C., Kiminami, C.S., “Directional and Rapid Solidification of Al-Nb-Ni Ternary Eutectic Alloy”, Mater. Sci. Eng. A, 375-377, 565–570 (2004) (Experimental, Interface Phenomena, 15) Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.10238.1.20, (2004) (Phase Diagram, Phase Relations, Review, 164) Costa, C.A.R., Batista, W.W., Rios, C.T., Milenkovic, S., Goncalves, M.C., Caram, R., “Eutectic Alloy Microstructure: Atomic Force Microscopy Analysis”, Appl. Surf. Sci., 240(1-4), 414–423 (2005) (Crys. Structure, Experimental, Morphology, 17) Belomyttsev, M.U., Laptev, A.I., Ezhov, I.P., Chertov, S.S., “Strength and Creep of Structural Materials Based on Intermetallic Compound NiAl”, Phys. Met. Metallogr. (Engl. Transl.), 101(5), 397–403 (2006), translated from Fiz. Metal. Metallov., 101(5), 429–435 (2006) (Experimental, Mechan. Prop., Morphology, 6) Chen, H., Du, Y., “Refinement of the Thermodynamic Odelling of the Nb-Ni System”, Calphad, 30, 308–315 (2006) (Calculation, Phase Diagram, #, *, 37) Greenberg, B.A., Antonova, O.V., Ivanov, M.A., Patselov, A.M., Plotnikov, A.V., “Some Features of the Formation and Destruction of Dislocation Barriers in Intermetallic Compounds: II. Observation of Blocked Superidislocations Upon Heating Without Stress”, Phys. Met. Metallogr. (Engl. Transl.), 102(1), 69–75 (2006), translated from Fiz. Metal. Metallov., 102(1), 749–755 (2006) (Crys. Structure, Experimental, Morphology, 5) Hagihara, K., Yokotani, N., Umakoshi, Y., “Temperature and Orientation Dependence of Fracture Behavior of Directionally Solidified Duplex-Phase Crystals Composed of Ni3X-Type Intermetallic Compounds”, Mater. Sci. Forum, 512, 67–72 (2006) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 12) Hu, R., Nash, P., “Review: Experimental Enthalpies of Formation of Compounds in Al-Ni-X Systems”, J. Mater. Sci., 41(3), 631–641 (2006) (Experimental, Thermodyn., 101) Raghavan, V., “Al-Nb-Ni (Aluminum-Niobium-Nickel)”, J. Phase Equilib. Diffus., 27(4), 397–402 (2006) (Crys. Structure, Phase Diagram, Review, 28)


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40

11 [2006Rin]

[2006Yu]

[2007Yu]

[Mas2] [V-C2]

Al–Nb–Ni Rinkevich, A.B., Stepanova, N.N., Burkhanov, A.M., “Acoustical Properties of Ni3Al Single Crystals Alloyed With Cobalt and Niobium”, Phys. Met. Metallogr. (Engl. Transl.), 102(6), 632–636 (2006), translated from Fiz. Metal. Metallov., 102(6), 678–682 (2006) (Experimental, Mechan. Prop., Phys. Prop., Thermodyn., 13) Yu, P., Kim, K.B., Das, J., Baier, F., Xu, W., Eckert, J., “Fabrication and Mechanical Properties of Ni-Nb Metallic Glass Particle-Reinforced Al-Based Metal Matrix Composite”, Scr. Mater., 54(8), 1445–1450 (2006) (Experimental, Mechan. Prop., Morphology, Phase Relations, 30) Yu, P., Zhang, L.C., Zhang, W.Y., Das, J., Kim, K.B., Eckert, J., “Interfacial Reaction During the Fabrication of Ni60Nb40 Metallic Glass Particles-Reinforced al Based MMCs”, Mater. Sci. Eng. A, 444(1-2), 206–213 (2007) (Experimental, Morphology, 31) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

DOI: 10.1007/978-3-540-88053-0_11 ß Springer 2009

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Al–Nb–Si

12

Aluminium – Niobium – Silicon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Christian Baetzner, Joerg Beuers, Michael Hoch, updated by Kostyantyn Korniyenko

Introduction Phase relations in the ternary Al-Nb-Si system are of great interest above all because alloys of this system are the splendid superconducting materials and also the perspective materials for application in the niobium silicide-based composites with the excellent practical properties similar to the nickel-based superalloys. Experimental data on phase equilibria in the Al-Nb-Si system are represented mainly by the series of isothermal sections at various temperatures [1961Bru, 1971Mue, 1972Pan, 1973All, 1977Mue, 1984Pan, 2001Mur1, 2003Zha]. Crystal structures of the phases taking part in equilibria in the Al-Nb-Si system are reported in [1961Bru, 1961Now, 1971Mue, 1974Joh, 1975Kha, 1977Ale, 1977Gur, 1977Mue, 1978Cat, 1978Dew, 1978Gol, 1987Ves, 2001Mur1, 2003Man, 2003Mur, 2003Zha, 2006Mat]. Experimental study of thermodynamic properties has been done in [1971Dub]. A thermodynamic assessment of the Al-Nb-Si system was carried out in [2004Sha] using the CALPHAD method. Two isothermal sections and the liquidus surface were calculated. Publications concerned with experimental studies of phase relations, crystal structures and thermodynamics, and applied techniques are listed in Table 1. Reviews on phase relations in the Al-Nb-Si system are presented in [1963Eng, 2005Rag, 2006Rag]. The character of phase equilibria was assessed in the previous MSIT evaluation by [1993Bae]. In comparison with that, the present report is supplemented by the information from later publications.

Binary Systems The Al-Si boundary binary system is accepted from the MSIT evaluation by [2004Luk]. The binary boundary Al-Nb and Nb-Si systems are based on [Mas2].

Solid Phases Crystallographic data on the known solid unary, binary and ternary Al-Nb-Si phases and their concentration and temperature ranges of stability are presented in Table 2. The τ1, Nb3Al2Si5, ternary phase has been detected by X-ray diffraction in pressed powder samples annealed at 1400˚C [1961Bru, 1961Now] and the structure and an X-ray pattern are given by [1961Now]. The phase equilibria above 25 at.% Nb were also studied by [1961Bru] and another phase, τ2, appears in the ternary which is suggested to be the Al-stabilized high temperature form of the βNb5Si3 [1961Now]. This view is supported by [1971Mue] who Landolt‐Bo¨rnstein New Series IV/11E1

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DOI: 10.1007/978-3-540-88053-0_12 ß Springer 2009

1

2

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Al–Nb–Si

concludes the extension of the Nb5Si3 phase up to 21 at.% Al at 1840˚C from metallography and X-ray diffraction in agreement with the binary βNb5Si3 Ð αNb5Si3 transition at 1650˚C [Mas2]. This view is also supported by [1984Pan] who found the βNb5Si3 phase in the ternary Nb5(AlxSi1–x)3 samples with x ranging from 0.6 to 0.75 at 1500˚C [1984Pan] and also at 1400˚C [1961Bru]. It was noted in the assessment of [1993Bae] that the solubility of 11 at.% Al in the αNb5Si3 phase directing towards NbAl3 at 1500˚C [1984Pan] is inconsistent with the solubility of less than 2 at.% Al at 1400˚C given by [1961Bru]. The ternary τ1 phase was also detected by [1973All] in a study of the liquid-solid equilibria at 1500 and 1300˚C (below 33 at.% Nb) by electromagnetic separation of phases and chemical analysis, metallography, electron microprobe, X-ray diffraction and DTA. At 1145 ± 10˚C the invariant transition type reaction L + τ1 Ð NbAl3 + NbSi2 was proposed by [1973All]. But it was concluded in [2003Mur] on the basis of X-ray diffraction patterns of the Nb3Al2Si5 matrix compacts annealed at 1000 and 750˚C that the above-mentioned reaction does not occur at 1145 ± 10˚C and that the τ1 phase is stable below this temperature. This conclusion is confirmed by the participation of the τ1 phase in the phase equilibria at 1000˚C established by [2003Zha]. The homogeneity range of the superconducting Nb3Al phase has been studied and is given in the ternary diagrams at 1840˚C [1971Mue], 1820˚C [1984Pan], 1700˚C [1972Pan], 1500˚C [1984Pan] and 1400˚C [1961Bru]. The solubility of Si was found to range between 3 to 9 at.% Si. Only [1974Joh] concluded a large solubility of about 70 mol% Nb3Si in Nb3Al from X-ray data and Tc (critical temperature) data in sputtered thin film samples annealed at 750˚C for 3 h. The lattice parameter of Nb3Al decreases with Si content [1971Mue, 1974Joh, 1975Kha, 1977Ale, 1977Mue, 1978Cat].

Invariant Equilibria Temperatures, reaction types and phase compositions relating to the invariant equilibria of the system are listed in Table 3. The reaction scheme is shown in Figs. 1a, 1b. The presented scheme is mainly based on the calculated liquidus surface [2004Sha]. It is amended in accordance with the accepted binary boundary systems. The existence of the quasibinary eutectic L Ð Nb5Si3 + NbAl3 proposed in [1984Pan] contradicts to the constitution of the isothermal section at 1500˚C [1973All] and to the calculated liquidus surface projection [2004Sha] and not accepted in the present evaluation.

Liquidus, Solidus and Solvus Surfaces The liquidus surface projection of the Al-Nb-Si is presented in Fig. 2 according to the calculations of [2004Sha] and the constitution of the accepted boundary binary systems. It was predicted that neither the τ1 nor the τ2 phases can melt congruently. Melting of alloys containing the τ2 phase will not occur until above 1475˚C, being consistent with the experimental data [1961Bru, 1984Pan, 2001Mur1]. It was noted by [2004Sha] that for the development of refractory composites based on Nb5Si3 and NbSi2, one should avoid phase fields such as the Nb5Si3 + τ1, so as to avoid problems arising from partial melting and hence liquidinduced fracture at high service temperatures. DOI: 10.1007/978-3-540-88053-0_12 ß Springer 2009

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The alloy Nb42.7Al55.7Si1.6 (at.%) containing the binary eutectic Nb2Al-NbAl3 exhibits a regular microstructure characterized by the eutectic cells in the as-cast condition. Ternary element (Si) additions modify the as-cast eutectic microstructure from cellular to dendritic [2000Rio].

Isothermal Sections The Nb rich part of the isothermal section at 1840˚C is presented in Fig. 3 on the basis of data reported by [1971Mue] with amendments according to the accepted here binary edge diagrams. The amendments concerned the extension of the homogeneity ranges of the intermediate phases. It was communicated by [1984Pan] that at 1820˚C the homogeneity range of the Nb3Al phase exists inside the ternary system without contiguity to the boundary binary Al-Nb system, but this assumption was not confirmed by the relevant experimental data. Figure 4 shows the isothermal section at 1500˚C compiled on the basis of [1973All] data for the Al-NbAl3-NbSi2-Si phase region, results of [1984Pan] for the Nb-Nb5Si3-NbAl3 range and the data on the homogeneity ranges discussed in the section “Solid Phases”. The isothermal section at 1400˚C was constructed by [1961Bru] (reproduced in the review [1963Eng]) but it contradicts to the Al-Si binary phase diagram because at this temperature aluminium should be liquid. However, only the solid phases were shown in equilibria. Later it was calculated by [2004Sha]. The calculated section agrees well with the experimental data of [1961Bru] in the Nb rich corner and with the later data of [2001Mur1] concerning phase equilibria with the participation of the τ1 and τ2 phases. It is presented in Fig. 5. A tentative isothermal section at 1300˚C is shown in Fig. 6. It was adopted from the data of [1973All] for the Al-NbAl3-NbSi2-Si phase region and the data on the homogeneity ranges discussed above. The isothermal section at 1000˚C was constructed from the experimental data obtained by [2003Zha] using a high-efficiency diffusion-multiple approach. The phase fields positions in the Nb-NbSi2-NbAl3 region were obtained from the tri-junction area of the diffusion multiple annealed at 1000˚C for 2000 h, at the same time three-phase equilibria at the Al-Si side were estimated from the high-temperature data of [1973All]. The of calculation of the isothermal section at 1000˚C by [2004Sha] agrees well with the experimental data of [2003Zha]. This section is presented in Fig. 7.

Thermodynamics A calorimetric investigation of heats of solution of silicon and aluminium in aluminothermal alloys was carried by [1971Dub]. The heat of dissolution was reported as 100.3 ± 12.5 kJ·mol–1 when the mass ratio Si:Nb was 0.25 and the content of Al varied in the interval from 1.9 to 7.7 mass% and when the mass ratio Si:Nb was 0.60 and the content of Al varied from 1.3 to 7.1 mass%. Thermodynamic assessment of the Al-Nb-Si system was carried out in [2004Sha] by evaluating the available equilibrium data, using the CALPHAD method. The Nb-Si diagram was reassessed to ensure good description of phase equilibria in the ternary system. The assessed binary phase diagram is nearly identical with that of [Mas2] accepted in the present evaluation. They differ only in the invariant temperatures by a few degrees. Thermodynamic Landolt‐Bo¨rnstein New Series IV/11E1

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models for the intermediate phases are based on crystal structures, as well as orientation and solubility ranges of single-phase fields in the experimental isothermal sections at different temperatures. The assessed isothermal sections at the temperatures of 1400 and 1000˚C agree well with the experimental data obtained on the ternary alloys by [1961Bru, 2001Mur1] and [2003Zha], respectively. The liquidus surface projection also was calculated.

Notes on Materials Properties and Applications The Al-Nb-Si alloys are of great practical interest due to the various aspects. In the first place, they possess excellent superconductive properties, with high values of the superconducting transition temperatures (Tc). Also these alloys are the constituents of the niobium silicidebased composites that show great promise for application as the next generation turbine airfoil materials with significantly higher operating temperatures than the current generation advanced nickel-based superalloys. The Nb-Si binary composites possess excellent creep strength but poor oxidation resistance and poor room temperature fracture toughness, at the same time alloying with Al can improve the oxidation resistance [2003Zha]. The applied experimental techniques and studied types of properties are listed in Table 4. It was established by [1971Mue] that the alloy Nb75.3Al21.5Si3.2 (in at.%) possesses a value of the superconducting transition temperature of 19.2 K. A similar value (18.3 K) was measured by [1977Ale] on the Nb75Al19.7Si5.3 specimen homogenized at 1650˚C for 5 h and then annealed at 700˚C for 250 h. For the alloys Nb3Si-Nb3Al alloys (with the aluminium content of 5 at.% and higher) annealed at 750˚C this parameter attains about 14.5 K [1974Joh]. It was shown that additions of Si decrease the superconducting transition temperature of the Nb3Al-based alloys [1974Joh, 1978Cat]. According to [1978Gol], the character of the concentration dependence of the Tc for the films of the Nb3Al-based alloys with silicon additions is smooth with a maximal value of 18.8 K corresponding to the composition of Nb3Al0.8Si0.2. Aiming at further improvement of high-field critical current density, Jc, in the Nb3Al conductors, Si addition was attempted by [2004Ban1]. The value of microhardness for the single-crystal NbAl3.3Si0.01 was reported in [1977Gur] as 4.51 ± 0.39 GPa. Its electrical resistivity was 6.4·10–11 Ω·m or 3.2·10–10 Ω·m at the temperatures of 77 K and 298 K (25˚C), respectively. The effect of mechanical properties of the Nb-based solid solution on toughness and strength of multiphase alloys in the Al-Nb-Si ternary system was studied by [1999Mur]. The toughness of (Nb) single-phase alloys and multiphase alloys estimated from SP energy increases with decreasing of the total (Si + Al) content in (Nb). It was reported by [2001Mur1, 2001Mur2] that the matrix compact corresponding to the composition of the τ1, Nb3Al2Si5 phase showed extremely good oxidation resistance at 1300˚C, although the compact showed poor oxidation resistance at 750˚C.

Miscellaneous Analysis of the ternary Al-Nb-Si alloys by atomic absorption spectrometry was carried out by [1973Mol]. For each element the influence of the acid solvent (HF or HNO3), the influence of two other components and the flame conditions were tested. The formation of the high Tc Nb3Al compound was unsuccessfully attempted by [1978Dew] using solid state diffusion from DOI: 10.1007/978-3-540-88053-0_12 ß Springer 2009

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ternary bronzes. The formation of the Nb(AlxSi1–x)2 phase on niobium dipped into the Al-Si liquid saturated by Si was investigated by [1998Nan]. The Nb(AlxSi1–x)2 consisted of submicron-order fine grains. A liquid phase consisting of Al-Si was observed at the grain boundary of the intermetallic layers and at the interface between the product and the refractory metals. This implied the formation of the intermetallics by solution-reprecipitation process. A new method for surface modification based on the arc surface alloying was proposed in [2003Mat]. Its feasibility was investigated performing NbAl3 coated on niobium. When tungsten arc was used to melt an aluminium plate placed on a Nb block, the niobium surface was also melted and a melt pool of Al-Nb binary alloy was formed on a niobium block. The melt pool solidified into niobium aluminides on the surface of the Nb block, forming a thick NbAl3 layer on the top surface of the coating layer. When an Al-Si alloy plate was used instead of the aluminium plate, a niobium alumino-silicide layer was formed on the niobium block. A new fabrication method for multifilamentary Nb/Al-Si precursors was developed by [2004Ban2]. The drawability of the composite wires has been significantly improved by using intermediately a reel-to-reel rapid heating and quenching (RHQ) technique. Unlike the wellknown RHQ process for the Nb3Al processing, this new RHQ process was specially performed at an early stage of the precursor processing to yield a fine microstructure of the core by solidifying rapidly only the Al alloy core from the molten state, without any reaction with the Nb matrix. In authors opinion, using this technique, not only the workability of the core but also the hardness balance between the matrix and the core can be improved, thereby making the subsequent restack and draw process easier. Based on the optimal process parameters, fabrication of multifilamentary Nb/Al-Si precursors with a piece length of 30 m was achieved as laboratory samples, reducing the matrix ratio.

. Table 1 Investigations of the Al-Nb-Si Phase Relations, Structures and Thermodynamics Reference [1961Bru]

Method / Experimental Technique Pressing; sintering; X-ray diffraction

Temperature / Composition / Phase Range Studied 1400˚C; the whole range of compositions

[1961Now] X-ray diffraction

The NbSi2-based and τ1 phases

[1971Dub]

Isothermal water calorimetry; aluminothermy; sintering

1.3 to 9.5 mass% Al

[1971Mue]

High-frequency melting; anode oxidation; X-ray diffraction (Debye-Scherrer technique)

1840˚C; 60–100 at.% Nb

[1972Pan]

Arc melting; annealing; metallography; X-ray 1700˚C; the Nb rich corner diffraction

[1973All]

High-frequency melting; electromagnetic separation; optical microscopy; EMPA

[1974Joh]

Sputtering simultaneously from two different 750˚C; 75 at.% Nb binary alloy electrodes; arc melting; pressing; vacuum annealing; anodization; X-ray diffraction; SEM; optical microscopy


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Temperature / Composition / Phase Range Studied

[1975Kha]

Arc melting; annealing; X-ray diffraction; metallography

Nb3Al-based alloys with additions up to 10 at.% Si

[1977Ale]

Arc melting; chemical analysis; annealing; X-ray diffraction

1650, 700˚C; Nb3Al-based alloy with addition of 5.3 at.% Si

[1977Gur]

Spontaneous crystallization; chemical analysis; X-ray diffraction

NbAl3.3Si0.01

[1977Mue]

Arc melting; annealing; metallography; X-ray 1750˚C, as-cast state; 75 at.% Nb diffraction; EMPA

[1978Cat]

Arc melting; rapid quenching; X-ray diffraction;

Rapid quenching from 2000˚C; Nb3Albased alloys with additions up to 15 at.% Si

[1978Dew]

Annealing; EMPA; X-ray diffraction

Nb3Al-based alloys with Si additions

[1978Gol]

Films precipitation; EMPA; X-ray diffraction

600–800˚C; 75 at.% Nb

[1984Pan]

Arc melting; annealing; metallography; X-ray 1820 and 1500˚C; the range diffraction; DTA Nb-Nb5Si3- NbAl3

[1987Ves]

Heating of the pressed workpieces; X-ray diffraction

The Nb3Al-Nb3Si section (up to 70 at.% Si)

[2000Rio]

Arc melting; ingots growth; optical microscopy; SEM; EDS

Nb42.7Al55.7Si1.6 (at.%)

[2001Mur1] Ball milling; sintering; X-ray diffraction; SEM; thermogravimetric analysis (TGA)

1200–1600˚C

[2001Mur2] Ball milling; sintering; X-ray diffraction; SEM; thermogravimetric analysis (TGA)

1200–1700˚C; Nb3Al2Si5

[2003Man]

Mechanical alloying; ball milling; X-ray diffraction; HRTEM; SAD

Nb40Al40Si20, Nb40Al30Si30

[2003Mat]

Arc surface alloying; SEM; EMPA

Aluminide coating on Nb

[2003Mur]

Ball milling; sintering; X-ray diffraction; SEM; thermogravimetric analysis (TGA); EDS

750 and 1000˚C; the τ1, Nb3Al2Si5 phase- containing matrix compacts

[2003Zha]

Diffusion-multiple approach; hot isostatic 1000˚C pressing (HIP); SEM; EMPA; electro- discharge machining (EDM)

[2006Mat]

Self-propagating high temperature synthesis 1100˚C; NbAl3 and Nb5Si3 phases containing alloys (SHS); spark plasma sintering (SPS); X-ray diffraction; SEM; EMPA; optical microscopy

[2007Qu]

Arc melting; EDM; back-scattered electron image (BEI); X-ray diffraction

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As-cast state; Nb78Al6Si16 (at.%)


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Phase/ Temperature Range [˚C] (Al) (I) < 660.452



cF4 Fm3m Cu

a = 404.96

hP2 P63/mmc Mg

a = 269.3 c = 439.8

cI2 Im3m W

a = 330.04

T = 25˚C [V-C2]

x = 0, 0 < y ≤ 0.015, T = 577˚C [2004Luk] y = 0, 0 < x · 6·10–4, T = 661.4˚C [Mas2]

NbxAl1–x–ySiy

(Al) (II) NbxAl1–x–ySiy (Nb) < 2469

T = 25˚C, p = 20.5 GPa [V-C2] x = 0, 0 < y ≲ 0.048, T ≈ 605˚C, p = 2.1 GPa [2004Luk] T = 25˚C [V-C2]

x = 0, 0 < y ≤ 0.035, T = 1920˚C [Mas2] x = 0, 0 < y ≤ 0.005, T = 1770˚C [Mas2] y = 0, 0 < x ≲ 0.215, T = 2060˚C [Mas2]

NbxAl1–x–ySiy

(Si) < 1414

cF8 a = 543.06 Fd3m C (diamond)

T = 25˚C [V-C2]

x = 0, 0 < y ≤ 1.8·10–5, T = 577˚C [2004Luk]

NbxAl1–x–ySiy


Comments/References

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Phase/ Temperature Range [˚C] Nb3Al < 2060

Pearson Symbol/ Space Group/ Prototype cP8 Pm3m Cr3Si

Nb3(Al1–xSix)

Nb2Al < 1940

tP30 P42/mnm CrFe


18.6 to 25 at.% Al [Mas2] Strukturbericht designation: A15 a = 518.6 [V-C2] a = 517.3 x = 0, T = 1840˚C [1971Mue] a = 518 x = 0.2, in the alloy sintered at 1400˚C [1961Bru] a = 517.3 x = 0.12, T = 1840˚C [1971Mue] a = 518.8 to 517 0 ≤ x ≤ 0.4, in the as-cast alloys [1975Kha] a = 518.9 to 517.3 0.08 ≤ x ≤ 0.4, in the alloys annealed at 900˚C for 14 d [1975Kha] a = 516.1 x = 0.212, subsequent annealing at 1650 and 700˚C [1977Ale] a = 517 0 ≤ x ≤ 0.5, in the alloys annealed at 1750˚C for 1 h [1977Mue] a = 518.4 to 516.25 0 ≤ x ≤ 0.6, in the alloys quenched from 2000˚C [1978Cat] a = 519 x = 0.2, in the film precipitated and annealed at 600–800˚C [1978Gol] a = 518.4 to 517.4 0 ≤ x ≤ 0.7 [1987Ves]

a = 994.3 c = 518.6 a = 988.2 to 994.2 c = 517.6 to 515.3

NbAl3 < 1680

tI8 I4/mmm TiAl3

a = 384 c = 857

Nb3Si 1980 - 1770

tP32 P42/n Ti3P

a = 1022.4 c = 518.9

βNb5Si3 (h) 2520 - 1650

tI32 I4/mcm W5Si3

αNb5Si3 (r) < 1940

tI32 I4/mcm Cr5B3

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Comments/References

30 to 42 at.% Al [Mas2]; labeled elsewhere as σ phase [V-C2] in the binary alloys with 25 to 35 at.% Al sintered at 1400˚C [1961Bru] 75 at.% Al [Mas2] [V-C2] 25 at.% Si [Mas2] [V-C2] 37.5 to 40.5 at.% Si [Mas2] [V-C2]

a = 1004.0 c = 508.1

37.5 to 38.5 at.% Si [Mas2] [V-C2]

a = 657.0 c = 1188.4

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Phase/ Temperature Range [˚C] NbSi2 < 1940

Pearson Symbol/ Space Group/ Prototype hP9 P6222 CrSi2

Nb(AlxSi1–x)2 * τ1, Nb3Al2Si5 Nb(AlxSi1–x)2

oF24 Fddd TiSi2

* τ2, Nb10Al3Si3 Nb5(AlxSi1–x)3

tI32 I4/mcm W5Si3


Comments/References 66.7 at.% Si [Mas2] Strukturbericht designation: C40 [V-C2]

a = 481.9 c = 659.2 a = 478.7 to 480 c = 658 to 663

0 ≤ x ≤ 0.15, T = 1400˚C [1961Bru, 1961Now] 0.3 ≤ x ≤ 0.375 [1973All] Strukturbericht designation: C54 in the Nb33Al20Si47 alloy [1961Now]

a = 838.6 b = 489.1 c = 877.6

Strukturbericht designation: D8m 0.45 ≤ x ≤ 0.60, T = 1400˚C [1961Bru]

a = 1014 to 1019 c = 507

. Table 3 Invariant Equilibria Composition* (at.%) Reaction

T [˚C]

Type

Phase

Al

Nb

Si

L Ð βNb5Si3 + Nb3Al

-

e1

L

21

73

6

L + Nb3Al Ð βNb5Si3 + (Nb)

2000

U1

L

16

76

8

L + βNb5Si3 Ð Nb3Si + (Nb)

-

U2

L

4

81

15

L + Nb3Al Ð Nb2Al + βNb5Si3

-

U3

L

35

63

2

L + βNb5Si3 Ð αNb5Si3

1800

p4

L

44

34 .5

21 .5

L + βNb5Si3 + NbSi2 Ð τ1

1750

P

L

28

32

40

L + βNb5Si3 Ð αNb5Si3 + τ1

1700

U4

L

37

30

33

L + Nb2Al + NbAl3 + βNb5Si3

-

D

L

57

42

1

L + βNb5Si3 Ð αNb5Si3 + NbAl3

-

U5

L

60

32

8

L + αNb5Si3 Ð τ1 + NbAl3

1490

U6

L

69

18

13

Note: * - values are estimated from the diagram


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. Table 4 Investigations of the Al-Nb-Si Materials Properties Reference


Type of Property

[1971Mue]

Anode oxidation, superconductivity tests Superconducting transition temperature, critical current density

[1972Pan]

Superconductivity tests

Superconducting transition temperature, curves of the superconductive transition

[1974Joh]

Anode oxidation, sputtering electrodes superconductivity tests

Superconducting transition temperature

[1975Kha]



[1977Ale]

Superconductivity tests, string magnetometer tests

Superconducting transition temperature, paramagnetic susceptibility

[1977Gur]

Microdurometry, electrical resistivity measurements

Microhardness, electrical resistivity

[1977Mue]



[1978Cat]



[1978Dew]

Superconductivity tests (inductive technique)


[1978Gol]

Superconductivity tests (resistive potentiometry)


[1980Mat]

Superconductivity tests (inductive NQR spectra; superconducting transition technique); magnetic and electrical tests; temperature; critical magnetic field; nuclear quadrupole resonance (NQR) electrical resistivity

[1987Ves]

Superconductivity tests (inductive technique); magnetic measurements

Superconducting transition temperature, curves of the superconductive transition; critical magnetic field dependences on temperature

[1989Kri]

Superconductivity tests (four-probe technique)

Volt-ampere characteristics

[1999Mur]

Small Punch (SP) toughness tests; tensile tests at room temperature; compression tests at high temperatures

Toughness; strength

[2001Mur1] Vickers microhardness tests; compression Microhardness; yield stress; oxidation tests; oxidation tests resistance

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Type of Property

[2001Mur2] Vickers microhardness tests; compression Microhardness; yield stress; oxidation tests; oxidation tests; Knoop indentation resistance; elastic modulus; fracture method; Vickers indentation toughness; thermal expansion coefficient microfracture method; thermal expansion tests [2004Ban1] Superconductivity tests [2006Mat]

Critical current density

Vickers hardness tests; ultrasonic method; Hardness; Young’s modulus; shear four-point bending tests; Archimedes modulus; Poisson’s ratio; density; method bending strength; fracture toughness


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. Fig. 1a Al-Nb-Si. Reaction scheme, part 1

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. Fig. 1b Al-Nb-Si. Reaction scheme, part 2

Al–Nb–Si


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. Fig. 2 Al-Nb-Si. Calculated liquidus surface projection

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. Fig. 3 Al-Nb-Si. Nb-rich part of the isothermal section at 1840˚C


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. Fig. 4 Al-Nb-Si. Isothermal section at 1500˚C

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. Fig. 6 Al-Nb-Si. Tentative isothermal section at 1300˚C

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References [1961Bru]

[1961Now] [1963Eng]

[1971Dub]

[1971Mue] [1972Pan]

[1973All]

[1973Mol]

[1974Joh] [1975Kha]

[1977Ale]

[1977Gur]

[1977Mue]

[1978Cat]

[1978Dew]

[1978Gol]

[1980Mat]

Brukl, C., Nowotny, H., Benesovsky, F., “Investigations in the Ternary Systems: V-Al-Si, Nb-Al-Si, CrAl-Si, Mo-Al-Si or Cr(Mo)-Al-Si” (in German), Monatsh. Chem., 92, 967–980 (1961) (Crys. Structure, Phase Diagram, Experimental, *, 20) Nowotny, H., Benesovsky, F., Brukl, C., “The Ternary: Niobium-Aluminium-Silicon” (in German), Monatsh. Chem., 92(1), 193–196 (1961) (Crys. Structure, Experimental, 14) English, J.J., “Binary and Ternary Phase Diagrams of Columbium, Molybdenum, Tantalum and Tungsten”, Defense Metals Information Center, Batelle Memorial Institute, Columbus 1, Ohio, 183, (99–2)-63 (1963) (Phase Diagram, Review, *, 1) Dubrovin, A.S., Gorelkin, O.S., Demidov, Yu.Ya., Chirkov, N.A., Kostylev, L.S., Kolesnikova, O.D., “Calorimetric Investigation of Heats Solution of Silicon and Aluminium in Aliminothermal Alloys” (in Russian), Metalloterm. Process. Khim. Met. Mater. Conference (1971), 121–130 (1971) (Thermodyn., Experimental, 11) Mueller, A., “Superconductivity in the A15-Phase in the Nb-Al-Si System” (in German), Z. Naturforsch., 26(A), 1035–1039 (1971) (Crys. Structure, Phase Relations, Experimental, Superconduct., *, 7) Pan, V.M., Latysheva, V.I., Sudovtsov, A.I., “Superconductivity of Niobium-Aluminium-Silicon Alloys”, Phys. Met. Metallogr. (Engl. Transl.), 33, 180–183 (1972), translated from Fiz. Met. Metalloved., 33(6), 1311–1313 (1972) (Phase Diagram, Phase Relations, Experimental, Superconduct., *, 3) Allibert, C., Wicker, A., Driole, J., Bonnier, E., “Study of the System Niobium-Aluminium-Silicon. I. Partial Isothermal Sections at 1500 and 1300˚C and Behaviour of the Phase Nb(Si, Al)2” (in French), J. Less-Common Met., 31(2), 221–228 (1973) (Morphology, Phase Diagram, Phase Relations, Experimental, *, 4) Molins, R., Garden, J., Bozon, H., Driole, J., “Study of the System Niobium-Aluminium-Silicon. II. Analysis of Ternary Alloys Niobium-Aluminum-Silicon by Atomic Absorption Spectrometry” (in French), J. Less-Common Met., 31(2), 229–237 (1973) (Morphology, Experimental, Kinetics, 7) Johnson, G.R., Douglass, D.H., “Superconductivity in New A-15 Niobium Alloys”, J. Low Temp. Phys., 14(5), 575–595 (1974) (Crys. Structure, Morphology, Experimental, Superconduct., 27) Khan, H.R., Raub, Ch.J., “Structure and Superconductivity of Ternary and Quaternary A15 Phases Based on Nb3Al” (in German), Metall, 29(7), 673–677 (1975) (Crys. Structure, Morphology, Experimental, Electronic Structure, Superconduct., 10) Alekseyevskiy, N.Ye., Ageyev, N.V., Shamray, V.F., “Superconductivity of Some Three-Component Solid Solutions Based on the Compound Nb3Al”, Fiz. Met. Metalloved., 43(1), 29–35 (1977), translated from Fiz. Met. Metalloved. (USSR), 43(1), 38–44 (1977), (Crys. Structure, Phase Relations, Experimental, Electronic Structure, Magn. Prop., Superconduct., 14) Gurin, V.N., Korsukova, M.M., Popov, V.E., Elizarova, O.V., Belousov, N.N., Kuz’ma, Yu.B., “Solid Solutions of B, C, Si in Aliminides of Transition Metals” in “Single-Crystals of Refractory and Rare Metals, Alloys and Compounds” (in Russian), Akad. Nauk SSSR, Nauka, Moscow, 39–42 (1977) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., 6) Mueller, P., “Superconductivity in Quasibinary Alloys of the A3B-Nb3Si Type with A15 Structure” (in German), Z. Metallkd., 68(6), 421–427 (1977) (Crys. Structure, Morphology, Phase Relations, Experimental, Superconduct., 26) Caton, R., Sweedler, A.R., “The Dependence of the Superconducting Transition Temperature on Silicon Concentration in the NbAlSi Ternary System”, J. Less-Common Met., 60, 91–100 (1978) (Crys. Structure, Morphology, Experimental, Superconduct., 10) Dew-Hughes, D., Luhmann, T.S., “The Thermodynamics of A15 Compound Formation by Diffusion from Ternary Bronzes”, J. Mater. Sci., 13, 1868–1876 (1978) (Crys. Structure, Phase Relations, Calculation, Experimental, Interface Phenomena, Superconduct., 41) Golovashkin, A.I., Levchenko, I.S., Lobanov, N.N., Motulevich, G.P., “Properties of Films of Ternary Superconducting Nb3(AlSi) Alloy” (in Russian), Fiz. Met. Metalloved., 46(1), 45–49 (1978) (Crys. Structure, Morphology, Experimental, Superconduct., 7) Matukhin, V.L., Safin, I.A., Shamray, V.F., “Nuclear Quadrupole Resonance 93Nb in the Ternary Nb3AlBased Phases” (in Russian), Fiz. Met. Metalloved., 50(3), 526–532 (1980) (Morphology, Experimental, Electronic Structure, Magn. Prop., Superconduct., 12)

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Al–Nb–Si [1984Pan]

[1987Ves]

[1989Kri] [1993Bae]

[1998Nan]

[1999Mur]

[2000Rio] [2001Mur1]

[2001Mur2]

[2003Man]

[2003Mat]

[2003Mur]

[2003Zha]

[2004Ban1]

[2004Ban2]

[2004Luk]

[2004Sha]

12

Pan, V.M., Latysheva, V.I., Kulik, O.G., Popov, A.G., Litvinenko, E.N., “The Nb-NbAl3-Nb5Si3 Phase Diagram”, Russ. Metall. (Engl. Transl.), (4), 233–235 (1984), translated from Izv. Akad. Nauk SSSR, Met., (4), 225–226 (1984) (Phase Diagram, Experimental, *, 6) Vesnin, Yu.I., Starikov, M.A., “On the Features of the Superconductivity of the Nb-Al-X Solid Solutions with the A15 Structure” (in Russian), Doklady Akademii Nauk SSSR, USSR, 296(1-3), 98–100 (1987) (Crys. Structure, Experimental, Magn. Prop., Superconduct., 11) Krivko, N.I., “Investigation of a Superconductor-Silicon Interface” (in Russian), Fizika Tverdogo Tela (USSR), 31(6), 225–230 (1989) (Morphology, Experimental, Interface Phenomena, Superconduct., 7) Baetzner, C., Beuers, J., Hoch, M., “Aluminium - Niobium - Silicon”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, Document ID: 10.16068.1.20, (1993) (Crys. Structure, Phase Diagram, Assessment, 13) Nanko, M., Takahashi, A., Ogura, T., Kitahara, A., Yanagihara, K., Maruyama, T., “Formation of Intermetallic Compounds on Refractory Metals in Aluminum-Silicon Liquid”, Materials Science and Engineering Serving Society. Proceedings of the Third Okinaga Symposium on Materials Science and Engineering Serving Society, Elsevier Science, Amsterdam, Netherlands, 299–302 (1998) (Morphology, Experimental, Interface Phenomena, 7) cited from abstract Murayama, Y., Hanada, S., “Effect of (Si + Al) Content in Nb Solid Solution on Mechanical Properties of Multiphase Nb-Si-Al Alloys” (in Japanese), J. Jpn. Inst. Met., 63(12), 1519–1526 (1999) (Morphology, Experimental, Mechan. Prop., 18) cited from abstract Rios, C.T., Milenkovic, S., Caram, R., “Directional Growth of Al-Nb-X Eutectic Alloys”, J. Cryst. Growth, 211, 466–470 (2000) (Phase Relations, Experimental, 8) Murakami, T., Sasaki, S., Ichikawa, K., Kitahara, A., “Microstructure, Mechanical Properties and Oxidation Behavior of Nb-Si-Al and Nb-Si-N Powder Compacts Prepared by Spark Plasma Sintering”, Intermetallics, 9(7), 621–627 (2001) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., *, 17) Murakami, T., Sasaki, S., Ichikawa, K., Kitahara, A., “Oxidation Resistance of Powder Compacts of the Nb-Si-Cr System and Nb3Si5Al2 Matrix Compacts Prepared by Spark Plasma Sintering”, Intermetallics, 9(7), 629–635 (2001) (Morphology, Phase Relations, Experimental, Kinetics, Mechan. Prop., Phys. Prop., *, 17) Manna, I., Chattopadhyay, P.P., Banhart, F., Fecht, H.-J., “Solid State Synthesis of Al-Based Amorphous and Nanocrystalline Al-Nb-Si and Al-Zr-Si Alloys”, Z. Metallkd., 94(7), 835–841 (2003) (Crys. Structure, Morphology, Phase Relations, Experimental, 24) Matsuura, K., Koyanagi, T., Ohmi, T., Kudoh, M., “Aluminide Coating on Niobium by Arc Surface Alloying”, Mater. Trans., JIM, 44(5), 861–865 (2003) (Morphology, Phase Relations, Experimental, Interface Phenomena, 14) Murakami, T., Sasaki, S., Ito, K., “Oxidation Behavior and Thermal Stability of Cr-Doped Nb(Si, Al)2 and Nb3Si5Al2 Matrix Compacts Prepared by Spark Plasma Sintering”, Intermetallics, 11(3), 269–278 (2003) (Crys. Structure, Morphology, Phase Relations, Experimental, Interface Phenomena, Kinetics, *, 28) Zhao, J.-C., Peluso, L.A., Jackson, M.R., Tan, L., “Phase Diagram of the Nb-Al-Si Ternary System”, J. Alloys Compd., 360, 183–188 (2003) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, Interface Phenomena, #, 36) Banno, N., Takeuchi, T., Kikuchi, A., Iijima, Y., Inoue, K., Yuyama, M., Wada, H., “First Trial to Fabricate Nb3(Al, Si) Multifilamentary Superconductors by Rapid-Heating and Quenching (RHQ) Process”, AIP Conference Proceedings (USA), 711(2), 515–522 (2004) (Morphology, Experimental, Magn. Prop., Superconduct., 7) cited from abstract Banno, N., Takeuchi, T., Kikuchi, A., Iijima, Y., Inoue, K., Yuyama, M., Wada, H., “Multifilamentary Nb/Al-Ge and Nb/Al-Si Precursor Fabrication Using the Intermediately Rapid Heating and Quenching Technique”, Superconduct. Sci. Techn., 17(3): 320–326 (2004) (Morphology, Experimental, 5) cited from abstract Lukas, H.L., Lebrun, N., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 20.14887.1.20, (2004) (Crys. Structure, Phase Diagram, Assessment, 29) Shao, G., “Thermodynamic Assessment of the Nb-Si-Al System”, Intermetallics, 12(6), 655–664 (2004) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, #, 38)


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[2006Rag] [2007Qu]

[Mas2] [V-C2]

Al–Nb–Si Raghavan, V., “Al-Nb-Si-Ti (Aluminum-Niobium-Silicon-Titanium)”, J. Phase Equilib. Diffus., 26(6), 638 (2005) (Phase Relations, Review, 6) Matsuura, K., Kata, D.B., Lis, J.T., Kudoh, M., “Grain Refinement and Improvement in Mechanical Properties of Nb-Al-Si Intermetallic Alloys”, ISIJ Int., 46(6), 875–879 (2006) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 7) Raghavan, V., “Al-Nb-Si (Aluminum-Niobium-Silicon)”, J. Phase Equilib. Diffus., 27(2), 163–165 (2006) (Crys. Structure, Phase Diagram, Phase Relations, Review, 9) Qu, S., Han, Y., Song, L., “Effects of Alloying Elements on Phase Stability in Nb-Si System Intermetallics Materials”, Intermetallics, 15(5-6), 810–813 (2007) (Morphology, Phase Relations, Experimental, 10) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

DOI: 10.1007/978-3-540-88053-0_12 ß Springer 2009

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Al–Ni–V

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Aluminium – Nickel – Vanadium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Frederick H. Hayes, Peter Rogl, Eberhard Schmid, updated by Viktor Kuznetsov

Introduction A large number of studies of the system was devoted to alloying behavior of V in the Ni aluminides, mainly γ’Ni3Al and βNiAl [1959Gua, 1977Pel, 1983Och, 1984Och1, 1984Och2, 1988Mor, 1989Hon, 1994Jia, 1997Pri, 2001Zap, 2004Ish]. Some studies of the system were also motivated by the possible influence of V on catalytic action of Al rich nickel aluminides [1971Mya, 1977Mya]. Data on the formation of intermetallics in the system are briefly reviewed in [1990Kum]. [2005Rag] summarized recent publications on phase equilibria. Phase equilibria are studied fairly well. There exist data for primary fields of crystallization of phases [1977Mya], for the liquidus, solidus [2001Miu] and solvus [1989Hon, 1991Mis] surfaces in the Ni rich region, as well as several complete [1977Mya] and partial [1991Cot, 1997Pri] isothermal sections. Two vertical sections NiAl3-VAl3 and Ni2Al3-V5Al8 were studied by [1971Mya], and the section between Ni3Al and VNi3 by [1984Gup]. The quasibinary eutectic that exists on the NiAl-V section, was studied by [1977Pel, 1991Kim, 2000Mil, 2002Mil]. Thermodynamic data are restricted to the low-temperature heat capacity data [1999Dar] and enthalpy of formation of liquid on the section V/Ni = 3/7, presented by [2005Sud]. Some mechanical properties of alloyed γ’Ni3Al and βNiAl were studied by [1985Ino, 1988Dar, 1991Hay, 1991Wu]. The kinetic simulation of ordering process in the alloy Ni-6.6Al-15.1V (at.%) at 800˚C was performed by [2001Par]. A series of works of chinese authors [2003Zha, 2004Zha, 2005Hou, 2005Li1, 2005Li2, 2005Li3, 2005Zha, 2006Li, 2007Li1, 2007Li2] who used phase field model is devoted to theoretical studies of kinetics of transformations between γ, γ’Ni3Al and VNi3 phases and the resulting morphology. Experimental investigations of phase relations, structures and thermodynamics of phases are summarized in Table 1.

Binary Systems The binary phase diagram for Al-V was taken from [2000Ric] and for Ni-V from [Mas2, 1982Smi]. However, it shall be noted that a recent reinvestigation of the Al-V binary system by [2000Ric] revealed a significantly lower peritectic formation temperatures of 1408˚C for V5Al8 and 1270˚C for VAl3 than those accepted by [Mas2] (1670˚C and 1360˚C, respectively). For the Al-Ni binary, the latest version [2004Sal] evaluated within the MSIT Binary Evaluation Program is accepted. Landolt‐Bo¨rnstein New Series IV/11E1

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Solid Phases V substitutes for Al to a considerable degree both in NiAl [1965Ram, 1977Mya, 1991Cot, 1997Pri] and in Ni3Al [1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas, 1989Hon]. The solubility on Ni in the vanadium aluminides is very small [1971Mya, 1981Yin]. No ternary phases have been definitely reported for this system. In particular, [1975Mar] found for the alloy of VNi2Al composition, annealed at 500˚C, a two-phase structure, though with the possible participation of a phase with the MnCu2Al (cF16) structure (X-ray data were inconclusive). A phase with this structure, but with a totally different lattice space (580.31 pm as compared with 633.0 pm after [1975Mar]) was reported by [1999Dar] in an alloy annealed at 800˚C. Moreover, [2004Ish] reports a second-order transition with the tricritical point between the phases with CsCl (cP2) and MnCu2Al (cF16) structures in the Al-Ni-Ti-V quaternary system. Nevertheless, for the ternary system the present data seem not to be conclusive, and the existence of this phase needs more definitive confirmation. [1984Lia] did not manage to prepare a VNi6Al phase (neither stable or metastable), although similar phases MNi6Al exist in the systems with Nb and Ta (metastable NbNi6Al and stable TaNi6Al). Data on the solid phases are listed in Table 2.

Invariant Equilibria A partial reaction scheme given in Fig. 1, which is complete for the Ni rich corner, is based largely on the fields of primary crystallization determined by [1977Mya] and on the findings of [1977Pel] who observed a quasibinary eutectic at 1360˚C and 40 at.% V on the NiAl to V section. The existence of the latter was confirmed also by [1991Cot, 2000Mil]. Between this maximum and the Ni-V binary edge there are three invariant reactions: two transition reactions, U1 and U2, and a ternary eutectic reaction at 1180˚C in which the Ni rich liquid containing 15 at.% Al and 32 at.% V is, on the basis of the solubility data of [1959Gua, 1984Gup], in equilibrium with the Ni rich solid solution containing 7 at.% Al, 34 at.% V, NiAl containing 21 at.% Al, 28 at.% V, and σ phase containing 5 at.% Al, 53 at.% V. At the Al rich side of the saddle point there is a sequence of transition reactions towards the Al rich corner. Two of them, U3 and U4 can be identified from [1977Mya]. In both Fig. 1 and Fig. 2 the σ/σ’transition of the Ni-V σ phase [Mas2] has been disregarded. Unfortunately, the data for the compositions of the phases are nearly completely absent, so the invariant equilibria cannot be tabulated.

Liquidus, Solidus and Solvus Surfaces Figure 2 gives the primary crystallization field boundaries based on [1977Mya, 1977Pel]. Figures 3 and 4 present the data of [2001Miu] showing the dependence of the γ liquidus and solidus temperatures on the variation of Al at the parametric V content, and on the variation of V at the parametric Al content, respectively. Figure 5 gives the (Ni)-solvus curves for the Ni rich corner from [1989Hon, 1991Mis]. Beyond the solubility limit the two-phase field γ+γ’ is entered (see also Fig. 8 and Fig. 9 below). In the V rich end of Fig. 5 the precipitation of the congruently transforming VNi3 is indicated. DOI: 10.1007/978-3-540-88053-0_13 ß Springer 2009

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Isothermal Sections A partial isothermal section at 1300˚C is given in Fig. 6 [1991Cot, 2005Rag]. A partial isothermal section at 1200˚C is presented in Fig. 7 [1997Pri, 2005Rag]. The isothermal sections at 1100 and 800˚C shown in Fig. 8 and Fig. 9, respectively, are based primarily on the results of [1977Mya] modified to be consistent with the binaries, to exhibit close agreement with the X-ray results of [1965Ram] concerning equilibria among (V), NiAl and the σ phase at 800˚C, and to be consistent with the results of [1989Hon] for the Ni rich corner at both temperatures. In addition, the boundaries and extension limit of γ’Ni3Al are accepted from [1983Och, 1984Gup]. The results of [1994Jia] for the equilibria between γ, γ’Ni3Al and βNiAl phases, which are presented in a tabular form, are reproduced in Tables 3 and 4.

Temperature – Composition Sections The section βNiAl-V contains quasibinary eutectic at 40 at.% V, 1360˚C [1977Pel, 1991Cot, 2000Mil]. However at lower temperatures the composition of phases is shifting away from the join, as may be seen from the isothermal section above, so the section cannot be quasibinary. The vertical sections VAl3-NiAl3 and V5Al8-Ni2Al3 are presented in Figs. 10 and 11, respectively. Both are taken from [1971Mya]. Figure 12 gives the 75 at.% Ni isopleth from [1989Hon] which contains a saddle point for the eutectoid decomposition of the (Ni) solid solution into VNi3 and Ni3Al on cooling.

Thermodynamics [1999Dar] performed low-temperature (3.2 to 10.3 K) measurements of the heat capacity of the VNi2Al phase. The results, when treated in the standard way (Cp(T ) = γelT + CD(θ/T)), give γel = 14.17±0.10 mJ·mol–1·K2, θD = 359±1.9 K. This equation is valid only below 7 K. It should be noted that the existence of this phase is questionable (see Solid Phases), so the result may really correspond to the βNi(Al0.5V0.5) phase. The enthalpies of formation of liquid along the section xV/xNi = 3/7 are presented on Fig. 13 taken from [2005Sud].

Notes on Materials Properties and Applications The possibility of plastifying of γ’Ni3Al and βNiAl by V additions was considered by [1985Ino, 1988Dar, 1991Wu] from both theoretical and experimental points of view. [1991Cot, 1991Kim, 1994Ben, 2000Mil, 2002Mil] studied morphology and strength of the directionally solidified quasibinary eutectic βNiAl-V. Creep behavior of off-stoichiometric binary and ternary γ’Ni3Al was studied by [1991Hay]. [1995Ino] measured the strength of an intermetallic phase with nanoinclusions of amorphous phase.


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Miscellaneous [1985Mor] used the system among the tests for their method of predicting solubility limits in the γ phase. [1988Mor] investigated static displacements of atoms in the alloyed γ’Ni3Al compound. [1994Ben] studied the evolution of morphology of phases in the near-eutectoid alloy Ni5Al20V (at.%) at 650 to 900˚C. [1997Kai, 2001Kai] discovered the formation of abrupt compositional changes in the βNiAl phase grown in diffusion pairs Ni0.8Al0.2+V (as well as several other third components). [2000Boz, 2002Boz] suggested a simple theoretical scheme for prediction of alloying behavior of third elements in βNiAl. [2001Sav] investigated experimentally the ordering kinetics in the alloy Ni75Al21V4 (at.%). [2001Son] performed ab initio analysis of site preference of V in βNiAl. [2001Ter] suggested usage of thermal conductivity measurements for determination of site preferences in γ’ Ni3Al phase. [2003Oza] tested hydrogen permeability of fcc Al-Ni-V alloys. [2005Cos] suggested using atomic force microscopy as a tool for studying the morphology of eutectics and used it for study of NiAl-V one. [2006Tan] investigated role of antiphase boundaries in kinetics of the L12→DO22 transformation.

. Table 1 Investigations of the Al-Ni-V Phase Relations, Structures and Thermodynamics Reference


Temperature/Composition/ Phase Range Studied

[1959Gua]

Microhardness measurement, XRD

Ni3Al + up to 8 at.% V, 1150˚C

[1965Ram]

XRD

25 to 70 at.% Al, 25 and 50 at.%V (6 compositions), 800˚C

[1971Mya]

Metallography, DTA, XRD, hardness and microhardness measurements

Sections NiAl3-VAl3 and Ni2Al3V5Al8, 600 to 1400˚C

[1975Mar]

XRD

VNi2Al composition, 500˚C

[1977Mya]

DTA, XRD, metallography, density, hardness and 800 to 1100˚C electrical resistivity measurements

[1977Pel]

DTA, SEM, directional solidification

Section NiAl-V (3 compositions)

[1981Yin]

XRD

Section VAl3-NiAl3, room temperature

[1983Och]

XRD

Ni3AlxV1–x, x = 1 to 0.4, 1000˚C

[1984Gup]

DSC, XRD

Ni3AlxV1–x, x = 0 to 1, 1000˚C

[1984Lia]

Melt quenching, XRD

VNi6Al composition Ni3Al+ up to 8 at.% V, 1000˚C

[1984Och1] XRD

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Temperature/Composition/ Phase Range Studied


[1989Hon]

DTA, metallography and SEM-EDX analysis

Up to 25 at.% V + Al, 1000 to 1320˚C

[1991Cot]

Metallography, DTA, TEM, SEM, EPMA

NiAl + up to 20 at.% V, 1300 to 1370˚C

[1991Mis]

DTA, metallography and SEM-EDX analysis

Up to 25 at.% V + Al, 1000 to 1320˚C

[1994Jia]

Diffusion pairs technique, metallography, EPMA 2 to 6, 13 and 31 mass% Al, 1 to 5 mass% V, 800 to 1300˚C

[1997Pri]

Metallography, microhardness measurement, EPMA

20 to 30 at.% Ni, 9.5 to 22.5 at.% Al, 1200˚C

[1999Dar]

Adiabatic calorimetry

VNi2Al composition, 1.2 to 10 K

[2001Miu]

DTA

0 to 16 at.% Al, 0 to 17 at.% V

[2005Sud]

Calorimetry

Liquid phase, V/Ni = 3/7, up to 16 at.% V




Comments/References

(Al) < 660.452

cF4 Fm 3m Cu

a = 404.96 a = 408.4

pure Al at 25˚C [Mas2] at 0.17 at.% V [V-C]

γ, (Ni) < 1455

cF4 Fm 3m Cu

a = 352.40

at 25˚C [Mas2]

(V) < 1910

cI2 Im 3m W

a = 302.40 a = 307.55

at 25˚C [Mas2] at 50 at.% Al [V-C2]

NiAl3 < 856

oP16 Pnma Fe3C

a = 661.3 ± 0.1 b = 736.7 ± 0.1 c = 481.1 ± 0.1

[2004Sal]

Ni2Al3 < 1138

hP5 P3ml Ni2Al3

a = 402.8 c = 489.1

[2004Sal]

βNiAl < 1651

cP2 Pm 3m CsCl

a = 288.72 ± 0.02 a = 287.98 ± 0.02

at 50 at.% Ni [2004Sal] at 54 at.% Ni [2004Sal]


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Ni5Al3 < 723

oC16 Cmmm Pt5Ga3

a = 753 b = 661 c = 376

at 63 at.% Ni [2004Sal]

γ’Ni3Al < 1372

cP4 Pm 3m AuCu3

a = 357.18 ± 0.02

[V-C2]

VAl3 < 1360

tI8 I4/mmm TiAl3

a = 378.0 c = 832.2

[V-C2]

V5Al8 < 1673

cI52 I 43m Cu5Zn8

a = 923.4 ± 0.5

[V-C2]

V4Al23 < 736

hP54 P63/mmc V4Al23

a = 769.28 c = 1704

[1989Mur]

VNi3 < 1045

tI8 I4/mmm TiAl3

a = 354.3 c = 720.2 a = 354.1 c = 721.8

at 23.44 at.% V [P]

VNi2 < 922

oI6 Immm MoPt2

-

[Mas2]

σ < 1280

tP30 P42/mnm σCrFe

a = 895.4 c = 463.5 a = 899.6 c = 465.3

at 57.5 at.% V [1982Smi]

cP8 Pm 3n Cr3Si

a = 471.2


V3Ni ≲900

Comments/References

at 25.60 at.% V [P]

at 63.2 at.% V [1982Smi] [1982Smi]

. Table 3 Equilibrium Compositions of the γ and γ’Ni3Al Phases and V Partition Coefficients [1994Jia] γ phase

Temperature [˚C]

γ’Ni3Al phase

V, at.%

Al, at.%

V, at.%

Al, at.%

Partition coefficient kVγ’/γ

1300

3.05

16.6

3.09

19.6

1.01

1200

0.45

17.7

0.89

22.3

1.98

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. Table 3 (continued) γ phase

Temperature [˚C]

γ’Ni3Al phase

V, at.%

Al, at.%

V, at.%

Al, at.%

Partition coefficient kVγ’/γ

0.54

15.3

0.83

21.1

1.53

3.38

12.6

4.27

18.5

1.26

1000

1.93

12.3

3.85

17.6

1.99

900

3.15

9.9

4.73

17.0

1.50

3.64

9.0

5.50

16.4

1.51

1.95

10.4

3.43

19.3

1.76

1100

800

. Table 4 Equilibrium Compositions of the γ’Ni3Al and βNiAl Phases and V Partition Coefficients [1994Jia] γ’Ni3Al phase

βNiAl phase

V, at.%

Al, at.%

V, at.%

Al, at.%

Partition coefficient kVβ/γ’

1300

2.17

24.0

0.98

32.4

2.21

1100

2.81

24.1

0.74

35.4

3.80

5.29

21.2

2.56

35.0

2.07

Temperature [˚C]


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. Fig. 1 Al-Ni-V. Partial reaction scheme

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. Fig. 2 Al-Ni-V. Liquidus surface projection


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Al–Ni–V

. Fig. 3 Al-Ni-V. Dependence of the γ liquidus and solidus temperatures on Al variation at the parametric V content

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. Fig. 4 Al-Ni-V. Dependence of the γ liquidus and solidus temperatures on V variation at the parametric Al content


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. Fig. 5 Al-Ni-V. The γ (Ni) solvus surface

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. Fig. 6 Al-Ni-V. Partial isothermal section at 1300˚C


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. Fig. 7 Al-Ni-V. Partial isothermal section at 1200˚C

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. Fig. 8 Al-Ni-V. Isothermal section at 1100˚C


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. Fig. 9 Al-Ni-V. Isothermal section at 800˚C

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. Fig. 10 Al-Ni-V. Isopleth at 75 at.% Al


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. Fig. 11 Al-Ni-V. Vertical section V5Al8-Ni2Al3

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. Fig. 12 Al-Ni-V. Isopleth at 75 at.% Ni


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. Fig. 13 Al-Ni-V. Enthalpy of formation of liquid along the section xV/xNi = 3/7

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References [1959Gua] [1965Ram]

[1971Mya]

[1975Mar]

[1977Mya]

[1977Pel]

[1981Yin]

[1982Smi] [1983Och] [1984Gup]

[1984Lia] [1984Och1]

[1984Och2] [1985Ino]

[1985Mis]

[1985Mor] [1985Nas] [1988Dar]

[1988Mor]

Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al (γ’ Phase)”, Trans. Metall. Soc. AIME, 215, 807–814 (1959) (Experimental, Phase Diagram, Phase Relations, 27) Raman, A., Schubert, K., “On the Crystal Structure of Some Alloy Phases Related to TiAl3. III. Investigations in Several T-Ni-Al and T-Cu-Al Systems” (in German), Z. Metallkd., 56, 99–104 (1965) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 14) Myasnikova, K.P., Ponomareva, L.F., Pryakhina, L.I., Marshakov, I.K., “Examination of Alloys of the Systems NiAl3-VAl3 and Ni2Al3-V5Al8”, Russ. Metall. (Engl. Transl.), (1), 126–128 (1971), translated from Izv. Akad. Nauk SSSR, Met., (1), 186–188 (1971) (Experimental, Phase Relations, #, 5) Marazza, R., Ferro, R., Rambaldi, G., “Some Phases in Ternary Alloys of Titanium, Zirconium and Hafnium, with a MgAgAs or AlCu2 Type Structure”, J. Less-Common Met., 39(2), 341–345 (1975) (Crys. Structure, Experimental, Morphology, 11) Myasnikova, K.P., Markiv, V.Ya., Pryakhina, L.I., Motrychuk, G.Yu., “Phase Equilibria in the V-Ni-Al System and Some Alloy Properties”, Russ. Metall. (Engl. Transl.), (3), 192–199 (1977), translated from Izv. Akad. Nauk SSSR, Met., (3), 222–229 (1977) (Experimental, Phase Diagram, Phase Relations, Phys. Prop., #, 12) Pellegrini, P.W., Hutta, J.J., “Investigations of Phase Relations and Eutectic Directional Solidification in NiAl-V Join”, J. Cryst. Growth, 42, 536–539 (1977) (Experimental, Morphology, Phase Relations, Phys. Prop., *, 3) Ying-Hong, Z., Jing-Qi, L., Jiang-Xuang, Z., Cheng, C.S., “A Room-Temperature Section of the Phase Diagram of TiAl3- VAl3-MAl3 of the System Alloys of Al-Ti-V-M (M = Ni, Fe)”, Acta Phys. Sin. (Chin. J. Phys.), 30, 972–975 (1981) (Phase Relations, Experimental, 9) Smith, J.F., Carlson, O.N., Nash, P.G., “The Ni-V (Nickel-Vanadium) System”, Bull. Alloy Phase Diagrams, 3, 342–348 (1982) (Phase Diagram, Review, Phase Relations, Crys. Structure, #, 35) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Additions”, Bull. P.M.E. (T.I.T.), 52, 1–17 (1983) Phase Relations, 7) Gupta, A., Horton, J.A., Liu, C.T., “Phase Formation and Stability in the Pseudobinary Ni3Al-Ni3VAlloy System”, Metals Soc. AIME, Conf: High-Temp. All., Bethesda, Maryland, USA, 115–123 (1984) (Phase Relations, Experimental, Crys. Structure, Morphology, 10) Liang, W.W., Standley, R., Nash, P., Skowron, M., “The Relative Stabilities of the Ni6AlX (X = V, Nb, Ta) Phases”, J. Mater. Sci. Lett., 3(3), 259–261 (1984) (Experimental, Phase Relations, 5) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni(γ), Ni3Al(γ’) and Ni3Ga(γ’) Solid Solutions”, Bull. Res. Lab. Precis. Machin. Electron., Tokyo Inst. Technol., 53, 15–28 (1984) (Crys. Structure, Experimental, 66) Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32(2), 289–298 (1984) (Experimental, Phase Diagram, Phase Relations, 90) Inoue, A., Masumoto, T., Tomioka, H., Yano, N., “Microstructure and Mechanical Properties of Ductile Intermetallic Compounds Produced by Melt Quenching”, Int. J. Rapid Solidification, 1, 115–142 (1985) (Morphology, Mechan. Prop., Review, 28) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni(γ), Ni3Al(γ’) and Ni3Ga(γ’) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33, 1161–1169 (1985) (Experimental, Crys. Structure, 64) Morinaga, M., Yukawa, N., Ezaki, H., Adachi, H., “Solid Solubilities in Nickel-Based F.C.C. Alloys”, Philos. Mag. A, 51(2), 247–252 (1985) (Phase Relations, Theory, 26) Nash, P., “Nickel-Base Intermetallics for High Temperature Alloy Design”, Mater. Res. Soc. Conf.: HighTemp. Ordered Intermet. Alloys, Boston, 423–427 (1985) (Review, Phase Diagram, Phase Relations, 15) Darolia, R., Lahrman, D.F., Field, R.D., Freeman, A.J., “Alloy Modeling and Experimental Correlation for Ductility Enhancement in NiAl”, High-Temperature Ordered Intermetallic Alloys III, Mater. Res. Soc. Symp. Proc., 133, 113–118 (1989) (Mechan. Prop., Experimental, Theory, 14) Morinaga, M., Sone, K., Kamimura, T., Ohtaka, K., Yukawa, N., “X-Ray Determination of Static Displacements of Atoms in Alloyed Ni3Al”, J. Appl. Crystallogr., 21, 41–46 (1988) (Crys. Structure, Experimental, 18)


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13 [1989Hon]

[1989Mur] [1990Kum] [1991Cot] [1991Hay]

[1991Kim]

[1991Mis]

[1991Wu]

[1994Ben]

[1994Jia]

[1995Ino]

[1997Kai]

[1997Pri] [1999Dar]

[2000Boz] [2000Mil]

[2000Ric]

[2001Kai]

[2001Miu]

[2001Par] [2001Sav]

Al–Ni–V Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of γ’ Solvus in Ni-Al-X Ternary Systems”, Mater. Res. Soc. Symp. Proc., 133, 429–440 (1989) (Experimental, Phase Diagram, Phase Relations, #, 35) Murray, J.L., “The Al-V (Aluminum-Vanadium) System”, Bull. Alloy Phase Diagrams, 10, 351–357 (1989) (Phase Diagram, Crys. Structure, Review, Phase Relations, #, 34) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35(6), 293–327 (1990) (Crys. Structure, Phase Diagram, Review, 158) Cotton, J.D., Kaufman, M.J., Noebe, R.D., “Constitution of Pseudobinary Hypoeutectic β-NiAl+α-V Alloys”, Scr. Metall. Mater., 25, 1827–1832 (1991) (Phase Relations, Mechan. Prop., #, 9) Hayashi, T., Shinoda, T., Mishima, Y., Suzuki, T., “Effect of Off-Stoichiometry on the Creep Behavior of Binary and Ternary Ni3Al”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 617–622 (1991) (Experimental, Mechan. Prop., 7) Kim, Y.D., Wayman, C.M., “Effect of Vanadium Transformation Behavior and Martensite Morphology in a Ni-Al Shape Memory Alloy”, Scr. Metall. Mater., 25(8), 1863–1868 (1991) (Abstract, Crys. Structure, 0) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the γ Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123–130 (1991) (Assessment, Experimental, Phase Diagram, Phase Relations, #, 5) Wu, Y.P., Sanchez, J.M., Tien, J.K., “Effect of APB Microsegregation on the Strength of Ni3Al with Ternary Additions”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 87–94 (1991) (Mechan. Prop., Calculation, 22) Bendersky, L.A., Biancaniello, F.S., Williams, M.E., “Evolution of the Two-Phase Microstructure L12 + D022 in Near-Eutectoid Ni3(Al,V) Alloy”, J. Mater. Res., 9(12), 3068–3082 (1994) (Morphology, Experimental, 16) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (γ (A1), γ’ (L12) and β (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473–485 (1994) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 25) Inoue, A., Kimura, H., Sasamori, K., Masumoto, T., “High Strength Al-V-M (M = Fe, Co or Ni) Alloys Containing High Volume Fraction of Nanoscale Amorphous Precipitates”, Mater. Trans. JIM, 36, 1219–1228 (1995) (Mechan. Prop., Experimental, 21) Kainuma, R., Ikenoya, H., Ohnuma, I., Ishida, K., “Pseudo-Interface Formation and Diffusion Behaviour in the B2 Phase Region of NiAl-Base Diffusion Couples”, Def. Diffus. Forum, 143-147, 425–430 (1997) (Crys. Structure, Experimental, Phase Relations, Phys. Prop., 10) Prima, S.B., Morozova, E.A., Bega, N.D., “Phase Equilibria in Vanadium-rich Alloys of the V-N-Al System”, Powder Metall. Met. Cer., 36(7-8), 390–393 (1997) (Experimental, Phase Relations, #, 6) da Rocha, F.S., Fraga, G.L.F., Brandao, D.E., da Silva, C.M., Gomes, A.A., “Specific Heat and Electronic Structure of Heusler Compounds Ni2TAl (T = Ti, Zr, Hf, V, Nb, Ta)”, Physica B (Amsterdam), 269, 154–162 (1999) (Crys. Structure, Experimental, 20) Bozzolo, G., Noebe, R.D., Honecy, F., “Modeling of Ternary Element Site Substitution in NiAl”, Intermetallics, 8, 7–18 (2000) (Crys. Structure, Review, 34) Milenkovic, S., Coelho, A.A., Caram, R., “Directional Solidification Processing of Eutectic Alloys in the Ni-Al-V System”, J. Cryst. Growth, 211(1-4), 485–490 (2000) (Crys. Structure, Experimental, Morphology, 13) Richter, K.W., Ipser, H., “The Al-V Phase Diagram between 0 and 50 Atomic Percent Vanadium”, Z. Metallkd., 91(5), 383–388 (2000) (Crystal Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 13) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of γ’/β Interface Morphologies in NiAl-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168–175 (2001) (Experimental, Phase Relations, Kinetics, 21) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: V, Nb and Ta) Ternary Systems”, J. Phase Equilib., 22, 345–351 (2001) (Experimental, Phase Diagram, Phase Relations, #, 9) Pareige, C., Blavette, D., “Simulation of the FCC -> FCC+L12+DO22 Kinetic Reaction”, Scr. Mater., 44 (2), 243–247 (2001) (Experimental, Kinetics, Phase Relations, 8) Savin, O.V., Stepanova, N.N., Akshentsev, Yu.N., Rodionov, D.P., “Ordering Kinetics in Ternary Ni3AlX Alloys”, Scr. Mater., 45(8), 883–888 (2001) (Crys. Structure, Electr. Prop., Experimental, Kinetics, 18)

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Al–Ni–V [2001Son] [2001Ter]

[2001Zap]

[2002Mil] [2002Boz] [2003Oza]

[2003Zha]

[2004Ish]

[2004Sal]

[2004Zha] [2005Cos]

[2005Hou]

[2005Li1]

[2005Li2] [2005Li3]

[2005Rag] [2005Sud]

[2005Zha]

[2006Li]

[2006Tan]

13

Song, Y., Guo, Z.X., Yang, R., Li, D., “First Principles Study of Site Substitution of Ternary Elements in NiAl”, Acta Mater., 49, 1647–1654 (2001) (Calculation, Crys. Structure, Electronic Structure, 17) Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurement”, J. Mater. Res., 16(8), 2314–2320 (2001) (Calculation, Crys. Structure, Experimental, Transport Phenomena, 63) Zapolsky, H., Pareige, C., Marteau, L., Blavette, D., Chen, L.Q., “Atom Probe Analyses and Numerical Calculation of Ternary Phase Diagram in Ni-Al-V System”, Calphad, 25(1), 125–134 (2001) (Calculation, Phase Relations, Thermodyn., 14) Milenkovic, S., Caram, R., “Microstructure of the Microalloyed NiAl-V Eutectics”, Mater. Lett., 55(12), 126–131 (2002) (Experimental, Kinetics, Morphology, 9) Bozzolo, G.H., Noebe, R.D., Amador, C., “Site Occupancy of Ternary Additions to B2 Alloys”, Intermetallics, 10, 149–159 (2002) (Crys. Structure, Review, 27) Ozaki, T., Zhang, Y., Komaki, M., Nishimura, Ch., “Hydrogen Permeation Characteristics of V-Ni-Al Alloys”, Int. J. Hydrogen Energy, 28, 1229–1235 (2003) (Crys. Structure, Electrochemistry, Experimental, 8) Zhao, Y.H., Hou, H., Xu, H., Wang, Y.X., Chen, Z, Sun, X.D., “Atomic Scale Computer Simulation for Early Precipitation Process of Ni75Al6V19 alloy”, J. Mater. Sci. Techn., 19(Suppl. 1), 17–19 (2003) (Kinetics, Morphology, Calculation, 6) Ishikawa, K., Ohnuma, I., Kainuma, R., Aoki, K., Ishida, K., “Phase Equilibria and Stability of HeuslerType Aluminides in the NiAl-Ni2AlTi-Ni2AlY (Y: V, Cr or Mn) Systems”, J. Alloys Compd., 367(1-2), 2–9 (2004) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 20) Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminium-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.10238.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 164) Zhao, Y.H., Chen, Z., Wang, Y.X., Lu, Y.L., “Atomic-Scale Computer Simulation for Early Precipitation Process of Ni75Al10V15 Alloy”, Progr. in Natural Sci., 14, 241–246 (2004) (Kinetics, Calculation, 16) Costa, C.A.R., Batista, W.W., Rios, C.T., Milenkovic, S., Goncalves, M.C., Caram, R., “Eutectic Alloy Microstructure: Atomic Force Microscopy Analysis”, Appl. Surf. Sci., 240, 414–423 (2005) (Morphology, Experimental, 17) Hou, H., Zhao, Y.H., Chen, Z., Xu, H., “Prediction for the Early Precipitation Process of Ni75AlxV25–x System with Lower Al Concentration by the Phase-Field Model”, Acta Metallurgica Sinica, 41, 695–702 (2005) (Kinetics, Calculation, 18) Li, Y.-S., Chen, Z., Wang, Y.-X., Lu, Y.-L., “Computer Simulation of γ’ and θ Phase Precipitation of NiAl-V Alloy Using Microscopic Phase-Field Method”, Trans. Nonferrous Met. Soc. China, 15, 57–63 (2005) (Kinetics, Morphology, Calculation, 20) Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., “Computer Simulation of the Interphase Boundary Evolution in Ni75AlxV25–x Alloy”, J. Mater. Sci. Techn., 21, 395–398 (2005) (Kinetics, Calculation, 20) Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., Zhang, J.J, “Microscopic Phase-Field Simulation for Nucleation Incubation Time of Ni75AlxV25–x Alloy”, J. Cent.-South Univ. Technol., 12, 635–640 (2005) (Kinetics, Calculation, 16) Raghavan, V., “Al-Ni-V (Aluminum-Nickel-Vanadium)”, J. Phase Equilib. Diffus., 26, 273–275 (2005) (Crys. Structure, Phase Diagram, Phase Relations, Review, 15) Sudavtsova, V.S., Makara, V.A., Kudin, V.G., “Part 3 (Alloys of Nickel and Tin, Methods of Modeling and Prognosis of Thermodynamic Properties), Ch. 6. Thermodynamic Properties and Phase Equilibria of Nickel Alloys” (in Ukrainian) in “Thermodynamics of Metallurgical and Welding Melts”, “Logos” Publ., Kiev, 2005 (Thermodyn., Review, 192) Zhao, Y.H., Ju, D.Y., Hou, H., “Atomic-Scale Computer Simulation of Mixture Precipitation Mechanism for Ni75AlxV25–x Alloy”, Mater. Sci. Forum, 475-479, 3115–3118 (2005) (Calculation, Kinetics, Phase Relations, 4) Li, Y., Chen, Z., Lu, Y., Wang, Y., Chu, Z., “Computer Simulation of Ordered Interphase Boundary Structure of Ni-Al-V Alloy Using Microscopic Phase-Field Method”, Rare Metal Mater. Eng., 35, 200–204 (2006) (Kinetics, Morphology, Calculation, 14) Tanimura, M., Koyama, Y., “The Role of Antiphase Boundaries in the Kinetic Process of the L12→D022 Structural Change of an Ni3Al0.45V0.50 Alloy”, Acta Mater., 54, 4385–4391 (2006) (Crys. Structure, Experimental, Kinetics, Morphology, 26)


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[2007Li2]

[Mas2] [V-C] [V-C2]

Al–Ni–V Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., Lai, Q.B., “Microscopic Phase-Field Simulation of Atomic Migration Characteristics in Ni75AlxV25–x Alloys”, Mater. Lett., 61, 974–978 (2007) (Calculation, Experimental, Phase Relations, 17) Li, Y.S., Chen, Z., Lu, Y.L., Wang, Y.X., “Coarsening Kinetics of Intermetallic Precipitates in Ni75AlxV25–x Alloys”, J. Mater. Res., 22, 61–67 (2007) (Experimental, Kinetics, Morphology, Phase Relations, Thermodyn., 30) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

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Al–O–Zr

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Aluminium – Oxygen – Zirconium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Mireille Harmelin, updated by Olga Fabrichnaya

Introduction The ZrO2-Al2O3 system is of interest because of the excellent mechanical properties of these composites widely used as structural materials [2002Boc], implant materials and cutting tools [2004Bas, 2004Ker, 2004Lee, 2006San1, 2006San2], protective and thermal resistant coatings [2006Por, 2008Che]. Laminated Al2O3/(Al2O3-ZrO2) composites exhibited good tribological behavior [2000Tar, 2003Tos]. Nanocomposite materials obtained in this system find widespread applications in structural field because of their excellent mechanical properties [2006Jia, 2006Ran, 2007Kon]. Mesoporous nanocrystalline composites Al2O3-(Y-stabilized) ZrO2 also used in catalysts, as well as in adsorption, separation and photoelectric devices [2005Che]. Fully stabilized ZrO2 (6.5–9 mol% Y2O3) in its cubic form is used as electrolyte in solid oxide fuel cell (SOFC). The Al2O3 is used to improve thermo-mechanical properties of this material [1995Nav, 2005Cho]. The Al2O3-ZrO2 eutectic composites are prospective materials for SOFC and oxygen sensors working at high temperatures due to the combination of ionic conductivity with improved mechanical properties and corrosion resistance. The ZrO2-Al2O3 system is a part of the high-order system CaO-Al2O3-SiO2-UO2-ZrO2 important for modeling of chemical interactions between nuclear reactor core debris and concrete [1990Rel, 1993Bal]. The Al2O3-ZrO2 thin films deposited on Si substrate are promising candidates for gate dielectric in metal-metal oxide semiconductors [2001Mor, 2003Zhu, 2005Biz, 2006Biz]. Reinforcement of Al matrix alloys by Al2O3 and Al3Zr by direct melt reaction improves mechanical properties of these materials [2001Zha, 2003Zha2, 2007Zha]. Most of the experimental studies on phase equilibria, microstructure development and properties are devoted to the ZrO2-Al2O3 system. Isothermal sections for the Al-O-Zr system were constructed by [1977Guk, 1978Guk1, 1978Guk2, 1978Guk3] based on diffusion couple studies. No ternary phases have been observed. The equilibrium quasibinary section ZrO2-Al2O3 has been determined by [1932War, 1964Alp, 1967Alp, 1994Lak, 1997Lak, 2000Jer, 2005Kam] and from 20 mass% Al2O3 by [1968Cev]. The quasibinary section ZrO2-Al2O3 was calculated by [1979Doe, 1980Wei, 1990Rel, 1992Wu, 1993Bal, 2004Fab, 2006Lak] using the CALPHAD approach. The previous MSIT evaluation of the experimental studies up to 1986 was presented by [1987Har]. Table 1 contains summary of experimental and theoretical studies from 1986 up to the present time.


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Al–O–Zr

Binary Systems The binary system Al-Zr is accepted according to the critical evaluation of [2004Sch]. The binary system O-Zr is accepted from [2006Wan]; information about high pressure phases is from [V-C2, 2005Oht, 1999Des]. The binary phase diagram of the Al-O system is accepted from [1992Tay]. Crystallographic data for solid phases are from [1985Wri] and from [1998Lev] and [V-C2] for metastable Al2O3 phases.

Solid Phases There are no ternary phases in the Al-O-Zr system. Monoclinic αZrO2 practically does not dissolve Al2O3. According to [1964Alp] the solubility of ZrO2 in Al2O3 is 0.83 mol%. [2000Jer] found that it is even smaller (0.008 mol% ZrO2). There is uncertainty in the solubility of Al2O3 in the tetragonal (βZrO2) and fluorite (γZrO2) structures of ZrO2. Maximal solubility of Al2O3 in the tetragonal phase βZrO2 is ranging from 1.3 to 7 mol% according to different experimental data: 7 mol% [1967Alp], 5 mol% [1997Lak] or 1.3 mol% [2000Jer]. [1997Lak] determined maximal solubility of Al2O3 in the fluorite phase γZrO2 to be 6 mol% based on thermal analysis data. [2006Lak] derived maximal solubility of Al2O3 in the tetragonal phase βZrO2 as 4.5 mol% and in the fluorite phase γZrO2 as 3 mol% by a Calphad type assessment. According to [1988Tan, 1994Bal, 1994Ish] non-equilibrium monoclinic, tetragonal and cubic phases could contain up to 20–40 mol% Al2O3. The tetragonal phase could have different ratio of a/c lattice parameters [1986Ady, 1988Tan, 1994Bal, 1994Ish]. A metastable tetragonal phase formed by diffusionless transformation on quick cooling is called t’ and its c/a ratio slightly exceed unity. In several works [1990Jay, 1991And, 1994Bal, 1995Nar, 2003Kin, 2007Pod], a formation of metastable Al2O3 phases, such as γ, δ or θ Al2O3, was observed during crystallization from amorphous and metastable solid solution of Zr(1–x)AlxO(2–x/2) or by rapid solidification. The γAl2O3 phase has a structure of spinel, the other phases δ and θ are the superstructures of γAl2O3. Investigation of structures of Al2O3 metastable phases was performed in [1998Lev] using TEM. A high temperature hexagonal phase referred to as the ε phase, was found by [1968Cev] above 1930˚C, at the Al2O3 side of the ZrO2-Al2O3 section, with a = 784.9 and c = 1618.3 pm. Crystallographic data for the solid phases are presented in Table 2.

Quasibinary Systems The ZrO2-Al2O3 quasibinary system was experimentally studied in several works [1932War, 1961Suz, 1964Alp, 1967Alp, 1968Cev, 1987Vol, 1997Lak, 2000Jer, 2005Kam]. Samples were melted in an Ar atmosphere by [1964Alp, 1967Alp] and in vacuum and tungsten crucibles by [1968Cev]. In all cases, temperature was measured with an optical pyrometer. [1994Lak, 1997Lak] studied liquidus in the ZrO2-Al2O3 system using derivative thermal analysis in air in solar furnace at the temperatures up to 3000˚C. Jerebtsov et al. [2000Jer] and Kamaev et al. [2005Kam] studied phase transformations in the ZrO2-Al2O3 system using DTA and XRD. There is a remarkable scatter in the eutectic temperature and composition obtained in the above mentioned experimental studies. The eutectic reaction was observed at 1920˚C by DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009

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[1932War], 1880˚C by [1964Alp, 1967Alp] and 1710±10˚C by [1968Cev]. The composition of the eutectic point was found to be 61.95 mol% Al2O3 by [1968Cev] and 63.93 mol% Al2O3 by [1964Alp, 1967Alp]. [1968Cev] observed no solid solution at low temperature between Al2O3 and ZrO2 on the Al2O3 side and [1964Alp, 1967Alp] found that less than 1 mass% ZrO2 goes into the solid solution in αAl2O3. According to [1964Alp, 1967Alp] about 7 mass% Al2O3 dissolves in solid ZrO2. Lakiza et al. [1994Lak, 1997Lak] determined a metatectic reaction between fluorite and tetragonal structures of ZrO2 and liquid at 2260˚C with the liquid composition of 20 mol% Al2O3 and the maximum solubility in ZrO2 of 5 mol%. The eutectic reaction was determined at 1860˚C and 63 mol% Al2O3 by [1994Lak, 1997Lak]. The results of [2000Jer] and [2005Kam] for the eutectic reaction are close to each other. The composition was obtained the same 64.47 mol% ZrO2 in both works and the temperature of 1866˚C [2000Jer] and 1861˚C [2005Kam]. The recommended phase diagram was derived by thermodynamic calculations taking into account phase equilibrium data obtained in the ZrO2-Al2O3 system together with the high-order system ZrO2-Al2O3-Y2O3 by [2006Lak]. The calculated phase diagram of the quasibinary system based on the thermodynamic description of [2006Lak] is presented in Fig. 1. It should be mentioned that the shape of the liquidus obtained in the works of [1964Alp, 2000Jer, 2005Kam] show a tendency to phase separation in the liquid state at temperatures slightly exceeding the liquidus line and compositions 55-45 mol% Al2O3.

Invariant Equilibria The ternary invariant equilibria of the quasibinary section ZrO2-Al2O3 are given in Table 3.

Isothermal Sections [1977Guk, 1978Guk1, 1978Guk2, 1978Guk3] have studied the reaction diffusion within the ternary system with metal/oxide diffusion couples annealed at the temperatures between 1000 and 1300˚C by means of electron microprobe analysis. Two isothermal sections at 1300 and 1130˚C have been deduced from the phases developing at the interface of the (Zr,Al)/ZrO2 and (Zr,Al)/Al2O3 samples. Figures 2 and 3 show the isothermal sections at 1300 and 1130˚C after [1978Guk2] with a minor correction due to the absence of Zr4Al3 above 1030˚C. In the diffusion couples Zr4Al3 has been observed and not the high temperature phase Zr5Al4 because of the transformations during cooling. It is interesting to note that a very small Al solubility in (αZr) was observed in the diffusion couples at the αZr/βZr interface [1978Guk1]. Since the two-phase field (αZr)+(βZr) extends from the O-Zr to the Al-Zr system (metastable above 940˚C), the tie lines of this field must shift eventually to end up with a higher Al content in (αZr) than in (βZr). From the binary O-Zr system it is presumable that αZrO2 occurs in the isothermal section at 1130˚C and βZrO2 at 1300˚C.

Thermodynamics A CALPHAD type assessment of thermodynamic parameters in the ZrO2-Al2O3 system was performed by [1979Doe, 1980Wei, 1990Rel, 1992Wu, 1993Bal, 2004Fab, 2006Lak]. Different Landolt‐Bo¨rnstein New Series IV/11E1

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models of liquid solution were used in the calculations. [1979Doe, 1980Wei] used the ideal solution model for the ZrO2-AlO1.5 liquid, [1992Wu] applied the quasi-chemical model. [1990Rel, 1993Bal] used the simple substitutional model with the excess Gibbs energy term expressed by the Redlich-Kister polynomial. No solubility in the ZrO2 phases was taken into account by [1979Doe, 1980Wei, 1992Wu, 1990Rel, 1993Bal]. Partially the ionic liquid model was used for the liquid phase and the compound energy formalism for the ZrO2-based solid solutions by [2004Fab, 2006Lak]. Excellent agreement with the experimental data of [1964Alp, 1967Alp] was obtained.

Notes on Materials Properties and Applications The ZrO2-Al2O3 composite materials are attractive due to their excellent mechanical properties combining properties of ZrO2 and Al2O3 [1997He, 2004Bas]. Introduction of ZrO2 in the Al2O3 matrix leads to composite materials with increased toughness due to the stress induced transformation of tetragonal to monoclinic ZrO2 phase [1997He, 2004Bas, 2006Ran]. Introducing Al2O3 in the ZrO2 matrix also improves thermal shock resistance and hardness in comparison with ZrO2 ceramics [1997He]. [1997He] studied mechanical properties (wear and friction test) of 20 mass% Al2O3 dispersed in the Y2O3 stabilized tetragonal ZrO2 matrix (Y-TZP) and Al2O3-15 mass% ZrO2 materials. [2004Bas] investigated mechanical properties of ZrO2(Y-TZP)-28 vol% Al2O3 composite. Influence of other metal oxides on mechanical properties of the Y-TZP-Al2O3 composite was studied by [2004Ker]. [2004Lee, 2005Lee, 2005Kim] produced Al2O3-m-ZrO2/t-ZrO2 fibrous composites using extrusion process and measured their mechanical properties. [2004Llo] studied influence of the Y2O3 on mechanical properties of directionally solidified Al2O3-ZrO2 eutectic. [2006San1, 2006San2] studied effect of composition and sintering condition on mechanical properties of ZrO2-Al2O3 biocompatible composites. Influence of sintering method (conventional one or in the arc plasma reactor) on density, mechanical and dielectric properties of the ZrO2-Al2O3 composites was studied by [2005Sah, 2006Sah]. [2000Tar, 2003Tos] prepared laminated composites by superimposing alternated layer of Al2O3/ZrO2 composite and ZrO2 and studied mechanical properties of these materials. The studied composites demonstrated improvement of surface toughness, reduced friction coefficient and increased wear resistance. [2002Cic] studied conductivity of the Al2O3-ZrO2 directionally solidified eutectic composites. These composites have good combination of mechanical and electrical properties making them prospective materials for SOFC and oxygen sensors working at high temperatures. [2005Cho] studied mechanical properties of Y2O3-stabilized cubic ZrO2 (10YSZ) reinforced with 0–30 mol% Al2O3. [2001Zha, 2003Zha2, 2007Zha] studied mechanical properties of Al matrix alloys reinforced by Al2O3 and Al3Zr produced by direct melt reaction. Tensile tests [2003Zha2] indicated that (Al3Zr+Al2O3)/Al-alloy composite exhibited high strength both at room and elevated temperatures. [2001Mor] studied stability of ZrAlxOy thin film obtained by sputtering on a Si substrate for a possible replacement of SiO2 as gate dielectric material in Si-based complimentary metaloxide semiconductors. [2003Zhu, 2005Biz, 2006Biz] measured dielectric constant of Al-O-Zr thin film deposited on a Si substrate as an alternative for gate dielectric application. Also Zr-doped Al2O3 is promising candidate for gate dielectric materials [2003Jun]. DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009

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Experimental studies of materials properties in the Al-O-Zr system are summarized in Table 4.

Miscellaneous [1984Kli] studied the rate of oxygen absorption by Al-Zr alloys (16.5, 16 mass% Al) at 100–500˚C, because they are used as gas absorber to protect various devises in electronic industries. [1993Pal] studied oxidation kinetics of Zr2Al3 in the stream of dry oxygen at atmospheric pressure with constant heating rate or isothermally at 602–692˚C. Activation energy was estimated from the obtained results to be 270 kJ·mol–1. [1994Dyb] suggested different mechanisms for oxidation process studied by [1993Pal]. [1994Pal] studied oxidation rate for αZr-1mass%Al, Zr3Al, Zr2Al, ZrAl, Zr2Al3, ZrAl2 and ZrAl3 with the increasing temperature and isothermally. ZrO2 in the tetragonal and monoclinic forms was found as main oxidation product, while Al escapes by diffusion in the bulk of alloy. [1985Gar1] and [1985Gar2] applied end-point thermodynamic analysis to tetragonal/ monoclinic transformation in the Al2O3-ZrO2 composites in the presence and absence of applied stress field. [1985Gar1] took into account the contribution of chemical, dilatational, residual shear strain and interfacial energy to total Gibbs energy and calculated dependence of transformation temperature from inversed critical size of particles. [1985Gar2] modelled microcrystal of tetragonal ZrO2 constrained in a matrix subjected to hydrostatic tensile stress field. [1995Nar] studied thin films grown on cubic YSZ by liquid precursor route. Phase transformation and epitaxy of these films were studied as a function of the heat treatment temperature and time. [1999Raj] developed a process of in-situ reduction of ZrO2 with excess aluminium for preparation of an Al-Zr master alloy. [2001Zha, 2001She, 2003Deq, 2003Zha1, 2007Zha] investigated the process of in-situ fabrication of Al matrix composite reinforced by Al3Zr and Al2O3. [2003Deq] used differential scanning calorimetry to study temperature of reduction of the ZrO2 by molten Al. [2002Pen] studied influence of physicochemical treatment and Al2O3 additions to transformation of the metastable tetragonal ZrO2 phase to monoclinic modification and showed that pressing at elevated temperature accelerates the formation of the equilibrium phase, whereas the Al2O3 additions stabilize a metastable phase. [2005Cal] studied effect of the solidification rate on the microstructure of the directionally solidified eutectic. [2004Muo] studied wetting and spreading behavior of the Ag-Cu-Ti alloys on aluminazirconia ceramics important for production of metal-ceramic joints.


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. Table 1 Investigations of the Al-O-Zr Phase Relations, Structures and Thermodynamics Reference



[1986Ady]

DTA, XRD, immersion method (refractive index determination)

300–1800˚C, ZrO2-Al2O3 (15–85 mol% ZrO2)

[1987Vol]

Melting, XRD, vibromilling, annealing

797–1747˚C, ZrO2-30 mass% Al2O3

[1988Tan]

Co-precipitation, heat treatment, cooling with different rates, XRD

1000, 1250˚C, ZrO2-Al2O3 0–21 mol% Al2O3

[1990Jay]

Co-precipitation, TEM, SEM/EDX, metallography, XRD

1797˚C sintering, ZrO2-Al2O3 (0–90 mass% ZrO2)

[1990Rel]

CALPHAD

1727–2727˚C, ZrO2-Al2O3

[1991And]

Rapid solidification by melt-extraction technique, TEM, SEM, XRD

1000–2000˚C, Al2O3-ZrO2 (25, 42 mol% ZrO2)

[1993Bal]

CALPHAD

1000–2800˚C, ZrO2-Al2O3

[1993Pal]

Arc-melting, XRD

602–692˚C, Zr2Al3 oxidation

[1994Bal]

Synthesis from pecursors, spray pyrolyzed 25–1200˚C (DTA), heat treatment and heat treated; 1000˚C, ZrO2-Al2O3 (10, 40 mol% XRD, DTA/TGA, electron diffraction, Raman Al2O3) spectroscopy, TEM, SEM

[1994Ish]

Co-precipitation, high isostatic pressing, heat treatment, XRD, TEM, SEM, DTA, IR spectroscopy

[1994Lak, 1997Lak]

DTA in He and TA in air in solar furnace, XRD 1100–3000˚C, ZrO2-Al2O3

[1994Pal]

XRD, EPMA, thermobalance

[1995Nar]

Thin film growth by liquid precursor route, 1100–1400˚C, ZrO2-20 mol% Al2O3 XRD, SEM, TEM

[1995Nav]

Co-precipitation, surface area 25–1000˚C, ZrO2-3–8 mol% Y2O3-5–20 measurements, XRD, TG/DTA, IR, SEM, TEM mass% Al2O3

[2000Jer]

DTA, X-ray fluorescence microanalysis

1800–2130˚C, ZrO2-Al2O3

[2001Mor]

Film sputtering on Si upon thermal annealing in vacuum or in O2, XRD, NRA, X-ray photoelectron spectroscopy, ion scattering spectroscopy

600˚C, Zr80Al20 in oxygen-containing plasma

[2001She]

TG/DTA, XRD

25–1100˚C, 55 vol% Al-25 vol% α-Al2O3 -20 vol% β-ZrO2

[2001Zha]

XRD, SEM, EPMA, TEM

800˚C, in-situ reaction ZrOCl2 with liquid Al (composite: Al matrix with Al3Zr and Al2O3)

DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009

1250˚C, ZrO2-Al2O3 (5–25 mol% Al2O3)

374–830˚C, Zr-1 mass% Al, Zr3Al, Zr2Al, ZrAl, Zr2Al3, ZrAl2, ZrAl3 oxidation

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[2002Cic]

XRD, light microscopy

1860˚C, 54–59 mol% Al2O3- 39–44 mol % ZrO2 (1–2 mol% Y2O3)

[2003Mpe]

Sintering, SEM

1500–1600˚C, 3YTZ (3 mass% Y2O3 ZrO2) - 0–75 vol% Al2O3

[2003Deq]

Uniaxial pressing, heating in Ar, DSC, XRD, SEM, thermodynamic calculations

150–1200˚C; Al-10–20 mass% ZrO2

[2003Jun]

X-ray emission and absorption spectroscopy

Zr-doped Al2O3

[2003Kin]

Plasma Al and Zr interaction with O2 gas, XRD, X-ray photoelectron spectroscopy (XPS), TEM/EDS

1000–1200˚C annealing, ZrO2-Al2O3

[2003Zha1] [2003Zha2]

Direct melt reaction, SEM, TEM, EPMA

800˚C, ZrOCl2+Al

[2003Zho]

Plasma spray method, sintering, annealing, 1400˚C, ZrO2(3 mol%Y2O3)-20, XRD, SEM 57 mass% Al2O3

[2003Zhu]

Laser deposition on Pt coated Si and Si substrate at 20 Pa O2, XRD, DTA

800–950˚C deposition, 25–1300˚C DTA; ZrO2-50 mol% Al2O3

[2004Fab]

CALPHAD

227–2727˚C, ZrO2-Al2O3

[2005Cal]

Arc-melting, XRD, SEM/EDX, micro-Raman

ZrO2-Al2O3 eutectic, rapid solidification

[2005Che]

XRD, TG/DSC, IR spectroscopy, TEM/EDS, nitrogen adsorption

100–1400˚C, ZrO2-Al2O3

[2005Kam]

DTA, XRD, SEM/EDX, calculations

1800–2100˚C; ZrO2-Al2O3

[2006Jia]

Coating (layer-by-layer method), cold and hot isostatic pressing, sintering, XRD, SEM, TEM

1350–1550˚C, Al2O3-12 mass% ZrO2

[2006Lak]

CALPHAD

1727–2727˚C, ZrO2-Al2O3

[2006Por]

Coating (magneton sputtering) in Ar and O2 gas mixture from metallic target, XRD

ZrO2-2–9 mass% Al2O3

[2006Ran]

Gel-precipitation, XRD, DSC/TG, Archimedes 25–1200˚C (DSC),1400–1600˚C method, IR-spectroscopy (sintering); Al2O3-ZrO2

[2006Sah] [2005Sah]

Conventional sintering and arc melting, XRD, SEM

1400–1600˚C, ZrO2-Al2O3

[2006San1] [2006San2]

Cold pressing, sintering, XRD, SEM

1500–1600˚C, ZrO2-0–30 mass% Al2O3

[2007Kon]

Pechini sol-gell process, HIP, DTA/TGA, XRD, 600–1500˚C, ZrO2-20 mass% Al2O3 TEM/EDX, SEM

[2007Pod]

Co-precipitation, heat treatment, XRD, DTA 700–1400˚C, ZrO2-50 mol% Al2O3


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. Table 1 (continued) Temperature/Composition/Phase Range Studied

Reference


[2007Sch]

Microwave plasma process, XRD, TEM, Perturbed Angular Correlation (PAC) measurement

≤627˚C, ZrO2 nanocrystalline particles covered by Al2O3

[2007Zha]

In-situ magneto-chemistry reaction, XRD, SEM, TEM

900˚C, Al-x mass% Zr(CO3)2; x = 5–25

[2008Car]

XRD, Neutron diffraction, stress 550 MPa

500˚C, 20 mass% ZrO2 - Al2O3

[2008Che]

Co-precipitation, plasma spray deposition, heat-treatment, XRD, SEM, chemical analysis (XPS – X-ray photoelectron spectroscopy)

1100–1500˚C, Al2O3-40 mass% ZrO2 (7 mass% Y2O3)





Comments/References

(αAl) < 660.452

cF4 Fm 3m Cu

a = 404.96

at 25˚C [Mas2]

(βAl)

hP2 P63/mmc Mg

a = 269.3 c = 439.8

25˚C, 20.5 GPa [Mas2]

(βZr) 1855 - 863

cI2 Im 3m W

a = 360.90

[Mas2]

(αZr) < 863

hP2 P63/mmc Mg

a = 323.16 c = 514.75

at 25˚C [Mas2]

Zr3Al < 1019

cP4 Pm 3m AuCu3

a = 437.2

[2004Sch]

Zr2Al(r) < 1250

hP6 P63/mmc Ni2In

a = 489.39 c = 592.83

[2004Sch]

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Comments/References

Zr5Al3(h) 1395 - 1000

tI32 I4/mcm W5Si3

a = 1104.4 c = 539.1

[2004Sch]

Zr5Al3(r)?

hP16 P6/mcm Mn5Si3

a = 817.0 c = 569.8

[2004Sch]

Zr3Al2 < 1480

tP20 P42/mnm Zr3Al2

a = 763.0 c = 699.8

[2004Sch]

Zr4Al3 ≤ 1030

hP7 P6/mmm Zr4Al3

a = 543.3 c = 539.0

[2004Sch]

Zr5Al4 1550 - 1000

hP18 P63/mcm Ti5Ga4

a = 844.7 c = 580.5

[2004Sch]

ZrAl < 1275

oC8 Cmcm CrB

a = 335.9 b = 1088.7 c = 427.4

[2004Sch]

Zr2Al3 < 1590

oF40 Fdd2 Zr2Al3

a = 960.1 b = 1390.6 c = 557.4

[2004Sch]

ZrAl2 < 1660

hP12 P63/mmc MgZn2

a = 528.24 c = 874.82

[2004Sch]

ZrAl3 < 1580

tI16 I4/mmm ZrAl3

a = 399.93 c = 1728.3

[2004Sch]

ZrAl3 (m)

cP4 Pm 3m AuCu3

a = 408

[2004Sch], metastable

αAl2O3 < 2054

hR30 R3c Al2O3

a = 475.4 c = 1299

[V-C2], corundum, congruent melting 2054˚C [Mas2]

γAl2O3

cF56 Fd 3m MgAl2O4

a = 794.7

[V-C2] metastable phase [Mas2]


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Comments/References

δAl2O3

oP64 P21212

a = 1589.4 b = 794.7 c = 1192.05

[1998Lev]

θAl2O3

mC8 C2/m Ga2O3

a = 1192.05 b = 280.97 c = 561.94 β = 104

[1998Lev]

λAl2O3

mP64 P21/c

a = 1685.81 b = 1589.4 c = 1192.05 β = 115

[1998Lev]

κAl2O3

hP44 P63mc Al2O3

a = 554.4 c = 902.4

[V-C2]

αZrO2 ≤ 1094˚C

mP12 P21/c ZrO2

a = 514.15 b = 520.56 c = 531.28 β = 99.30 a = 516.2 b = 519.4 c = 532.5 β = 99.11

baddeleyite [2006Wan]

a = 360.55 c = 517.97

up to 4.5 mol% Al2O3 for x = 0, at 1393˚C [V-C2] for 16.7 mol% Al2O3, Metastable [1988Tan] 5–40 mol% Al2O3 metastable t’ [1994Ish]

βZrO2 2311 - 1094

γZrO2 2710 - 2311

tP6 P42/nmc ZrO2

cF12 Fm3m CaF2

a = 359.5 c = 519.1 a = 508.5 c = 512.7 to 518.2 a = 513.2 a = 507 to 510.5

δZrO2

oP16 Pbcm ZrO2

DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009

a = 503.64 b = 525.46 c = 508.55

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for 16.3 mol% Al2O3, Metastable [1988Tan]

fluorite up to 3 mol% Al2O3 in equilibrium conditions, 5–40 mol% Al2O3 metastable [1994Ish] high pressure phase 3 < p < 12.5 GPa [2005Oht] for x = 0 [V-C2]


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Phase/ Temperature Range [˚C] εZrO2

oP12 Pnma PbCl2

ωZrO2

tP12 P4m2


Comments/References high pressure phase [1999Des] p > 12.5 GPa [2005Oht]

a = 562 b = 334.7 c = 650.3 a = 504.6 c = 521.9

high pressure phase 3.5 < p < 15 GPa [V-C2]


T [˚C]

Type

Phase

γZrO2 Ð L + βZrO2

2314

e1

γZrO2 βZrO2

L Ð βZrO2 + Al2O3

1856

e2

e3

Zr

1.9815

66.3356

31.6828

1.6832

66.3861

31.9307

L

14.2800

64.2867

21.4334

29.2985

61.7836

8.9179

2.9521

66.1746

30.8732

Al2O3 1130

O

L βZrO2

γZrO2 Ð αZrO2 + Al2O3

Al

40

60

0.0

γZrO2

0.878

66.5205

32.6015

αffl ZrO2

0.082

66.6530

33.2646

Al2O3

40

60

0.0

. Table 4 Investigations of the Al-O-Zr Materials Properties Reference


Type of Property

[1997He]

Friction and wear test (ball-on-plate wear Mechanical properties test)

[2000Tar]

Vickers hardness, indentation technique, wear test (pin-on-disc apparatus), XRD, SEM, EDAX

Mechanical properties (hardness, residual stress, Young’s modulus, wear resistance)

[2001Zha]

Tensile test

Mechanical properties


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. Table 4 (continued) Reference [2002Cic]

Method / Experimental Technique Impedance analysis for conductivity measurements

Type of Property Electric properties

[2003Mpe] Conductivity, dielectric constant

Electric properties

[2003Gan]

Compression test, SEM, optical microscopy

Mechanical properties of amorphous ZrO2 - 40 mol% Al2O3

[2003Tos]

Indentation technique, wear test (pin-on- Mechanical properties (residual stress, disc apparatus) surface toughness, friction coefficient, wear resistance)

[2003Zha2] Tensile test

Mechanical properties of (Al-7Si-0.3Mg)/ (Al3Zr+Al2O3)

[2004Bas]

Archimedes method, resonance frequency method, hardness test, indentation XRD, SEM, EPMA/EDS

Density, mechanical properties (elastic modulus, Vickers hardness, fracture toughness)

[2004Ker]

Wear friction test (pin-on-disc apparatus), Mechanical properties (friction XRD, SEM/EDX coefficient, wear resistance)

[2004Lee] [2005Lee] [2005Kim]

Archimedes method, hardness test, indentation, four point bending test, XRD, SEM, TEM

Density, mechanical properties (Vickers hardness, fracture toughness, bending strength)

[2004Llo]

Three point bending test, hardness test, indentation fracture method, XRD, SEM, TEM, Raman spectroscopy

Mechanical properties of directionally solidified eutectic Al2O3-ZrO2

[2005Cho]

XRD, SEM, TEM/EDS, mass-volume method, flexure strength test, fracture toughness test, impulse excitation of vibration method, microhardness test, fatigue test

Density and mechanical properties (fracture toughness, flexure strength, elastic properties, Vickers hardness, fatigue) ZrO2 (10 mol% Y2O3)-0–30 mol% Al2O3

[2006Biz] [2005Biz]

XRD, IR-spectroscopy, X-ray spectroscopy ZrO2-Al2O3 thin films, dielectric properties (XPS), capacitance measurements

[2006Sah] [2005Sah]

Archimedes method, Vickers indentation Density, mechanical properties method, dielectric bridge, XDT, SEM (hardness), dielectric properties ZrO2Al2O3 composites

[2006San1] Vickers indentation method [2006San2]

Mechanical properties (Vickers hardness, fracture toughness) of ZrO2-0–30 mass% Al2O3

[2007Zha]

Mechanical properties (tensile and yield strength, elongations) of Al/(Al3Zr+Al2O3) composite

Tensile test

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. Fig. 1 Al-O-Zr. Quasibinary system ZrO2-AlO1.5


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. Fig. 2 Al-O-Zr. Isothermal section at 1300˚C

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. Fig. 3 Al-O-Zr. Isothermal section at 1130˚C


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Tananaev, I.V., Khodzhamberdiev, M.S., Salyn, A.L., Beresnev, E.N., Pakhomov, V.I., Orlovskii, V.P., “Solubility of Al2O3 in ZrO2”, Inorg. Mater. (Engl. Trans.), 24, 1651–1653 (1988), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 24(11), 1925–1927 (1988) (Crys. Structure, Experimental, 10) Relave, O., Chevalier, P.Y., Cheynet, B., Cenerino, G., “Thermodynamic Calculation of Phase Equilibria in a Quinary Oxide System of First Interest in Nuclear Energy Field UO2-ZrO2-SiO2-CaO-Al2O3”, User Aspects of Phase Diagrams, Proc. Int. Conf., 25–27 June 1990, Petten, Hayes, F.H. (Ed.), Int. Met., 55–63 (1991) (Calculation, Phase Diagram, Phase Relations, 19) Ando, T., Shiohara, Y., “Metastable Alumina Structures in Melt-Extracted Alumina-25% Zirconia and Alumina-42% Zirconia Ceramics”, J. Am. Ceram. Soc., 74(1), 410–417 (1991) (Experimental, Morphology, Phase Diagram, Phys. Prop., Thermodyn., 23) Jayaram, V., Levi, C.G., Witney, T., Mehrabian, R., “Characterization of Al2O3-ZrO2 Powders Produced by Electrohydrodynamic Atomization”, Mater. Sci. Eng., 124(A), 65–81 (1990) (Experimental, Phase Relations, 15) Taylor, J.R., Dinsdale, A.T., Hillert, M., Selleby, M., “A Critical Assessment of Thermodynamic and Phase Diagram Data for the Al-O System”, Calphad, 16(2), 173–179 (1992) (Calculation, Phase Diagram, Thermodyn., 22) Wu, P., “Optimization and Calculations of Thermodynamic Properties and Phase Diagrams of Multicomponent Oxide Systems”, Ph. D. Thesis, Ecole Polytechnique, Montreal, Canada (1992) Ball, R.G.J., Mignanelli, M.A., Barry, T.I., Gisby, J.A., “The Calculation of Phase Equilibria of Oxide Core-Concrete Systems”, J. Nucl. Mater., 201, 238–249 (1993) (Thermodyn., Calculation, 66) Paljevic, M., “Selectiv Oxidation of Zr2Al3”, J. Alloys Compd., 191, 27–29 (1993) (Experimental, Interface Phenomena, Phase Relations, 13) Balmer, M.L., Lange, F.F., Levi, C.G., “Metastable Phase Selection and Partitioning for Zr(1–x) AlxO(2–x/2) Materials Synthesized with Liquid Precursors”, J. Am. Ceram. Soc., 77(8) 2069–2075 (1994) (Experimental, Phase Relations, 20) Dybkov, V.I., “Comment on the Paper (Selective Oxidation of Zr2Al3) by M. Paljevic”, J. Alloys Compd., 215, L1-L1 (1994) (Experimental, 3) Ishida, K., Hirota, K., Yamaguchi, O., “Formation of Zirconia Solid Solutions Containing Alumina Prepared by New Preparation Method”, J. Am. Ceram. Soc., 77(5), 1391–1395 (1994) (Crys. Structure, Experimental, Morphology, Phase Diagram, Phase Relations, 20) Lakiza, S.N., Lopato, L.M., Shevchenko, A.V., “Interaction in the Al2O3-ZrO2-Y2O3 System”, Powder Metall. Met. Ceram., 33(9-10), 486–490 (1994), translated from Poroshk. Metall., (9–10), 46–51 (1994) (Experimental, Morphology, Phase Diagram, Phase Relations, 22) Paljevic, M., “High-Temperature Oxidation Behavior in the Zr-Al System”, J. Alloys Compd., 204, 119–126 (1994) (Experimental, 57) Navarro, L.M., Recio, P., Duran, P., “Preparation and Properties Evaluation of Zirconia-Based Al2O3 Composites as Electrolytes for Solid Oxide Fuel Cell Systems. 1. Powder Preparation and Characterization”, J. Mater. Sci., 30(8), 1931–1938 (1995) (Experimental, Mechan. Prop., Morphology, Thermodyn., 19) Narwankar, P.K., Speck, J.S., Lange, F.F., “Phase Partitioning and Epitaxy of Zr(Al)O2 Thin Films on Cubic Zirconia Substrates”, J. Mater. Res., 10(7), 1756–1763 (1995) (Experimental, Morphology, Phase Relations, Phys. Prop., 22) He, Y.J., Winnubst, A.J.A., Burggraaf, A.J., Verweij, H., van der Varst, P.G.T., de With, G., “Sliding Wear of ZrO2-Al2O3 Composite Ceramics”, J. Eur. Ceram. Soc., 1371–1380 (1997) (Experimental, 25) Lakiza, S.M., Lopato, L.M., “Stable and Metastable Phase Relations in the System Alumina-ZircomiaYttria”, J. Am. Ceram. Soc., 80(4), 893–902 (1997) (Experimental, Phase Relations, 26) Levin, I., Gemming, Th., Brandon, D.G., “Some Metastable Polymorphs and Transient Stages of Transformation in Alumina”, Phys. Status Solidi A, 166(1), 197–218 (1998) (Crys. Structure, Experimental, Phase Relations, 24) Desgreniers, S., Lagarec, K., “High-Density ZrO2 and HfO2: Crystalline Structures and Equations of State”, Phys. Rev. B, 59(13), 8467–8472 (1999) (Crys. Structure, Experimental, 25) Rajagopalan, P.K., Sharma, I.G., Krishnan, T.S., “Production of Al-Zr Master Alloy Starting from ZrO2”, J. Alloys Compd., 285(1-2), 212–215 (1999) (Experimental, Thermodyn., 8) Jerebtsov, D.A., Mikhailov, G.G., Sverdina, S.V., “Phase Diagram of the System: Al2O3-ZrO2”, Ceram. Int., 26, 821–823 (2000) (Experimental, Phase Diagram, 10)


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Al–O–Zr Tarlazzi, A., Roncari, E., Pinasco, P., Guicciardi, S., Melandri, C., de Portu, G., “Tribological Behaviour of Al2O3/ZrO2-ZrO2 Laminated Composites”, Wear, 244(1-2), 29–40 (2000) (Experimental, Mechan. Prop., Morphology, Phys. Prop., 23) Morais, J., da Rosa, E.B.O., Pezzi, R.P., Miotti, L., Baumvol, I.J.R., “Composition, Atomic Transport, and Chemical Stability of ZrAlxOy Ultrathin Films Deposited on Si(001)”, Appl. Phys. Lett., 79(13), 1998–2000 (2001) (Experimental, 18) Sheedy, P.M., Caram, H.S., Chan, H.M., Harmer, M.P., “Effects of Zirconium Oxide on the Reaction Bonding of Aluminum Oxide”, J. Am. Ceram. Soc., 84(5), 986–990 (2001) (Experimental, 16) Zhao, Y. T., Sun, G. X., “In-Situ Synthesis of Novel Composites in the System Al-Zr-O”, J. Mater. Sci. Lett., 20(20), 1859–1861 (2001) (Crys. Structure, Experimental, Mechan. Prop.,6) Bocanegra-Bernal, M.H., Diaz De La Torre, S., “Review. Phase Transitions in Zirconium Dioxide and Related Materials for High Performance Engineering Ceramics”, J. Mater. Sci., 37(23), 4947–4971 (2002) (Crys. Structure, Mechan. Prop., Phase Relations, Review, Thermodyn., 175) Cicka, R., Trnovcova, V., Starostin, M.Y., “Electrical Properties of Alumina-Zirconia Eutectic Composites”, Solid State Ionics, 148(3-4), 425–429 (2002) (Experimental, Morphology, 14) Pentin, I.V., Oleinikov, N.N., Murav’eva, G.P., Eliseev, A.A., Tret’yakov, Yu.D., “Stability of Tetragonal ZrO2 Toward External Influences”, Inorg. Mater. (Engl. Trans.), 38(10), 1012–1014 (2002) (Crys. Structure, Experimental, 8) Deqing, W., Neumann, J., Lopez, H.F., “Reactive Synthesis of in-Situ ZrAl3-Al2O3-Al Composites”, Metall. Mater. Trans. A, 34(6), 1357–1360 (2003) (Crys. Structure, Experimental, Morphology, Thermodyn., 21) Gandhi, A.S., Jayaram, V., “Plastically Deforming Amorphous ZrO2-Al2O3”, Acta Mater., 51(6), 1641–1649 (2003) (Experimental, Mechan. Prop., 47) Jung, R., Lee, J.-Ch., So, Y.-W., Noh, T.-W., Oh., S.-J., Lee, J.-Ch., Shin, H.-J., “Bandgap States in Transition-Metal (Sc, Y, Zr, and Nb)-Doped Al2O3”, Appl. Phys. Lett., 83(25), 5226–5228 (2003) (Electronic Structure, Experimental, 7) Kinemuchi, Y., Mouri, H., Suzuki, T., Suematsu, H., Jiang, W., Yatsui, K., “Increase in Phase Transition Temperature of Activated Alumina with Nano-Zirconia Synthesized by Pulsed Wire Discharge”, J. Am. Ceram. Soc., 86(9), 1522–1526 (2003) (Crys. Structure, Experimental, Phase Relations, 16) M’Peko, J.-C., Spavieri Jr., D.L., da Silva, Ch.L., Fortulan, C.A., de Souza, D.P.F., de Souza, M.F., “Electrical Properties of Zirconia-Alumina Composites”, Solid State Ionics, 156(1-2), 59–69 (2002) (Electr. Prop., Experimental, 23) Toschi, F., Melandri, C., Pinasco, P., Roncari, E., Guicciardi, S., de Portu, G., “Influence of Residual Stresses on the Wear Behavior of Alumina/Alumina-Zirconia Laminated Composites”, J. Am. Ceram. Soc., 86(9), 1547–1553 (2003) (Calculation, Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 33) Zhao, Y.T., Cheng, X.N., Dai, Q.X., Cai, L., Sun, G.X., “Crystal Morphology and Growth Mechanism of Reinforcements Synthesized by Direct Melt Reaction in the System Al-Zr-O”, Mater. Sci. Eng. A, 360(1-2), 315–318 (2003) (Experimental, Morphology, Phys. Prop., 10) Zhao, Y.T., Dai, Q.X., Cheng, X.N., Sun, S.C., “Microstructure and Properties of In-Situ Synthesized (Al3Zr + Al2O3)p/A356 Composites”, Internat. J. Mod. Phys. B, 17(8-9), 1292–1296 (2003) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 10) Zhou, X., Shukla, V., Cannon, W.R., Kear, B.H., “Metastable Phase Formation in Plasma-Sprayed ZrO2 (Y2O3)-Al2O3”, J. Am. Ceram. Soc., 86(8), 1415–1420 (2003) (Crys. Structure, Phase Diagram, 13) Zhu, J., Liu, Z.G., “Structure and Dielectric Properties of Zr-Al-O Thin Films Prepared by Pulsed Laser Deposition”, Microelectr. Eng., 66(1-4), 849–854 (2003) (Crys. Structure, Experimental, Phys. Prop., 10) Basu, B., Vleugels, J., van der Biest, O., “ZrO2-Al2O3 Composites With Tailored Toughness”, J. Alloys Compd., 372, 278–284 (2004) (Experimental, Mechan. Prop., Morphology, 23) Fabrichnaya, O., Aldinger, F., “Assessment of Thermodynamic Parameters in the System ZrO2-Y2O3Al2O3”, Z. Metallkd., 95(1), 27–39 (2004) (Assessment, Phase Diagram, Thermodyn., 61) Kerkwijk, B., Garcia, M., van Zyl, W.E., Winnubst, L., Mulder, E.J., Schipper, D.J., Verweij, H., “Friction Behaviour of Solid Oxide Lubricants as Second Phase in α-Al2O3 and Stabilised ZrO2 Composites”, Wear, 256(1-2), 182–189 (2004) (Experimental, Morphology, Phys. Prop., 38) Lee, B.-T., Kim, K.-H., Han, J.-K., “Microstructures and Aterial Properties of Fibrous Al2O3-(mZrO2)/t-ZrO2 Composites Fabricated by a Fibrous Monolithic Process”, J. Mater. Res., 19(11), 3234–3241 (2004) (Crys. Structure, Experimental, Mechan. Prop., Morphology, Phys. Prop., 17)

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LLorca, J., Pastor, J.Y., Poza, P., Pena, J.I., de Francisco, I., Larrea, A., Orera, V.M., “Influence of the Y2O3 Content and Temperature on the Mechanical Properties of Melt-Grown Al2O3-ZrO2 Eutectics”, J. Am. Ceram. Soc., 87(4), 633–639 (2004) (Experimental, Mechan. Prop., Morphology, Phase Relations, 27) Muolo, M.L., Ferrera, E., Morbelli, L., Passerone, A., “Wetting, Spreading and Joining in the AluminaZirconi-Inconel 738 System”, Scr. Mater., 50, 325–330 (2004) (Experimental, Kinetics, Morphology, 31) Schuster, J.C., “Al-Zr (Aluminium - Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.13524.1.20, (2004) (Phase Diagram, Phase Relations, Assessment, 138) Bizarro, M., Alonso, J.C., Ortiz, A., “The Effect of the Process Conditions on the Synthesis of Zirconium-Aluminum Oxide Thin Films Prepared by Ultrasonic Spray Pyrolysis”, J. Electrochem. Soc., 152(11), F179-F184 (2005) (Experimental, Thermodyn., 29) Calderon-Moreno, J.M., Yoshimura, M., “Stabilization of Zirconia Lamellae in Rapidly Solidified Alumina-Zirconia Eutectic Composites”, J. Eur. Ceram. Soc., 25(8), 1369–1372 (2005) (Experimental, Morphology, 22) Chen, H., Wang, X.M., Shia, J., Xiao, P., Yan, D., “A Novel Structural Mesoporous Alumina/Yttrium Doped Zirconia Nanocrystalline Composite Derived by Solvothermal Approach”, J. Mater. Res., 20(1), 42–47 (2005) (Crys. Structure, Experimental, Morphology, Phase Relations, Thermodyn., 19) Choi, S.R., Bansal, N.P., “Mechanical Behavior of Zirconia/Alumina Composites”, Cer. I., 31(1), 39–46 (2005) (Mechan. Prop., 29) Kamaev, D.N., Archugov, S.A., Mikhailov, G.G., “Behavior of the Al2O3-ZrO2 System at High Temperatures”, Russ. J. Appl. Chem., 78(3), 347–350 (2005) (Experimental, Phase Diagram, Phase Relations, 16) Kim, T.-S., Goto, T., Lee, B.-T., “Microstructural Control and Mechanical Properties of Fibrous Al2O3/ ZrO2 Composites Fabricated by Extrusion Process”, Scr. Mater., 52(8), 725–729 (2005) (Experimental, Mechan. Prop., Morphology, Phase Relations, Phys. Prop., 10) Lee, B.-T., Jang, D.-H., Kang, I.-Ch., Lee, Ch.-W., “Relationship Between Microstructures and Material Properties of Novel Fibrous Al2O3-(m-ZrO2)/t-ZrO2 Composites”, J. Am. Ceram. Soc., 88(10), 2874–2878 (2005) (Experimental, Mechan. Prop., Morphology, Phase Relations, 21) Ohtaka, O., Andrault, D., Bouvier, P., Schultz, E., Mezouar, M., “Phase Relations and Equation of State of ZrO2 to 100 GPa”, J. Appl. Crystallogr., 38(5), 727–733 (2005) (Crys. Structure, Experimental, Phys. Prop., Thermodyn., 38) Sahu, D.R., Mishra, D.K., Swain, D., Ray, M., Roul, B.K., “Effect of Compositional Variation in Sintering Behaviour of Al-Zr Oxide Composites”, Mater. Sci. Eng. B, 119(1), 29–35 (2005) (Experimental, Morphology, Phase Relations, Phys. Prop., 32) Bizarro, M., Alonso, J.C., Fandino, J., Ortiz, A., “Effect of Water Addition in the Spray Solution on the Synthesis of Zr-Al Oxide Films Prepared by the Pyrosol Process”, J. Electrochem. Soc., 153(7), F153F159 (2006) (Experimental, 44) Jia, Y., Hotta, Y., Sato, K., Watari, K., “Homogeneous ZrO2-Al2O3 Composite Prepared by Nano-ZrO2 Particle Multilayer-Coated Al2O3 Particles”, J. Am. Ceram. Soc., 89(3), 1103–1106 (2006) (Experimental, Morphology, Phase Relations, 26) Lakiza, S., Fabrichnaya, O., Zinkevich, M., Aldinger, F., “On the Phase Relations in the ZrO2-YO1.5AlO1.5 System”, J. Alloys Compd., 420(1-2), 237–245 (2006) (Assessment, Crys. Structure, Experimental, Phase Diagram, Phase Relations, Thermodyn., 23) Portinha, A., Teixeira, V., Carneiro, J., Newton, R., Fonseca, H., “Structural Characterization of Sputtered Composite Stabilized ZrO2 Thin Films”, Mater. Sci. Forum, 514-516, 1150–1154 (2006) (Crys. Structure, Experimental, Nano, 18) Rana, R.P., Pratihar, S.K., Bhattacharyya, S., “Effect of Powder Treatment on the Crystallization Behaviour and Phase Evolution of Al2O3-High ZrO2 Nanocomposites”, J. Mater. Sci., 41(21), 7025–7032 (2006) (Crys. Structure, Experimental, 14) Sahu, D.R., Roul, B.K., Singh, S.K., Choudhury, R.N.P., “Studies on Sintering Behaviour of Al2O3ZrO2 Oxide Composites Processed by Extended Arc Thermal Plasma and Conventional Heating”, J. Mater. Sci., 41(17), 5480–5489 (2006) (Crys. Structure, Electr. Prop., Experimental, Morphology, 41) Santos, C., Teixeira, L.H.P., Daguano, J.K.M.F., Strecker, K., Elias, C.N., “Effect of Isothermal Sintering Time on the Properties of the Ceramic Composite ZrO2-Al2O3”, Mater. Sci. Forum, 530-531, 526–531 (2006) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 8)


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[2008Car]

[Mas2] [V-C2]

Al–O–Zr Santos, C., Teixeira, L.H.P., Strecker, K., Elias, C.N., “Effect of Al2O3 Addition on the Mechanical Properties of Biocompatible ZrO2-Al2O3 Composites”, Mater. Sci. Forum, 530-531, 575–580 (2006) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 8) Wang C., Zinkevich M., Aldinger F., “The Zirconia - Hafnia: DTA measurements and Thermodynamic Calculations”, J. Am. Ceram. Soc., 89(12) 3751–3758 (2006) (Experimental, Calculation, Phase Relations, Thermodyn., 78) Kong, Y.-M., Kim, H.-E., Kim, H.-W., “Production of Aluminum-Zirconium Oxide Hybridized Nanopowder and Its Nanocomposite”, J. Am. Ceram. Soc., 90(1), 298–302 (2007) (Experimental, Mechan. Prop., Morphology, Nano, 17) Podzorova, L.I., Ilicheva, A.A., Shvorneva, L.I., “Effect of the Precipitation Sequence on Phase Formation in the ZrO2-CeO2-Al2O3 System”, Inorg. Mater. (Engl. Trans.), 43(9), 972–975 (2007), translated from Neorg. Mater., 43(9), 1086–1089 (2007) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 14) Schlabach, S., Szabo, D.V., Vollath, D., de la Presa, P., Forker, M., “Zirconia and Titania Nanoparticles Studied by Electric Hyperfine Interactions, XRD and TEM”, J. Alloys Compd., 434-435, 590–593 (2007) (Crys. Structure, Experimental, Nano, 13) Zhang, S.-L., Zhao, Y.-T., Chen, G., Cheng, X.-N., Dai, Q.-X., “Microstructures and Mechanical Properties of Aluminum Matrix Composites Fabricated from Al-x wt.% Zr(CO3)2 (x = 5, 10, 15, 20, 25) Systems”, J. Alloys Compd., 429(1-2), 198–203 (2007) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 25) Chen, D., Jordan, E.H., Gell, M., Ma, X., “Dense Alumina-Zirconia Coatings Using the Solution Precursor Plasma Spray Process”, J. Am. Ceram. Soc., 91(1), 359–365 (2008) (Crys. Structure, Experimental, Mechan. Prop., Morphology, 20) Carter, G.A., van Riessen, A., “Neutron Diffraction of Zirconia-Dispersed Alumina with Increasing Stress and Temperature”, J. Am. Ceram. Soc., 91(2), 559–562 (2008) (Crys. Structure, Experimental, Mechan. Prop., 10) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

DOI: 10.1007/978-3-540-88053-0_14 ß Springer 2009

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Aluminium – Tantalum – Titanium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Tamara Velikanova, Mikhail Turchanin, Svitlana Ilyenko, Guenter Effenberg, Vasyl Tomashik, Dmytro Pavlyuchkov

Introduction Alloys based on the Al-Ta-Ti system are promising candidate materials for various structural applications. Alloys based on the γαffl phase (TiAl) have an excellent potential to become an important next generation aerospace materials because of their low density, high melting temperature, good elevated-temperature strength, high resistance to oxidation and hydrogen absorption, and excellent creep properties. The alloys of the Al-Ta-Ti system with the composition in the range 45–50 at.% Al are potential high temperature engineering materials. Additions of Ta to the γαffl phase (TiAl) are useful to improve the fracture toughness, ductility and corrosion and oxidation resistance at high temperatures. However, there is still a scope for alloy development to optimize microstructures of these alloys. Several experimental investigations of phase relations in the system were undertaken and, as a result, a set of isothermal sections at 1000 [1966Ram, 2000Kai], 1100 [1983Sri, 1991McC2, 1991Das, 1992Per, 1993Das, 1993Jew], 1200 [2000Kai], 1300 [2000Kai], 1330 [1991Boe], 1350 [1995Wea], 1440 [1991McC2, 1993Das], and 1450˚C [1991Wea, 1995Wea] is reported. Based on the solidification behavior of Al-Ta-Ti alloys [1991McC1, 1991McC2, 1992McC] constructed a partial liquidus surface. Table 1 summarizes the experimental studies on phase equilibria. The above mentioned experimental results were reviewed in [2005Das, 2005Rag].

Binary Systems The Al-Ta phase diagram from [1996Du] (Fig. 1), Ta-Ti from [Mas2] are accepted. For the AlTi system, the MSIT evaluation from [2004Sch] is available. More recently, a new review by [2006Sch] and a thermodynamic assessment by [2007Wit] appeared which show differences in two regions. [2006Sch] and [2007Wit] show two peritectoid reactions involving (αTi), (βTi) and Ti3Al and a small two-phase field between (βTi) and Ti3Al in the 1150–1200˚C temperature range. Evaluating the same experimental literature information, [2004Sch] concluded that Ti3Al transforms congruently and that there is no invariant equilibrium involving these three phases. According to [2004Sch], the different phase equilibria may be related to very small differences in Gibbs energy, i.e. very small driving forces, leading to phase transitions in that range. Thus, it is very difficult to decide which interpretation of the experimental data is actually correct.


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The second difference considers the equilibria between γ(TiAl), TiAl2 and the onedimensional antiphase domain structures or “long period structures (LP)” which is stable at high temperature in the range of 65 to 75 at.% Al. The assessments by [2006Sch] and [2007Wit] are based on more and recent experimental data and give a more complete discussion of the stable equilibria in this region. Thus the phase diagram calculated by [2007Wit] is accepted in the present evaluation because of its larger base of experimental data. The majority of these were assessed in [2006Sch] and complemented by own key experiments. The Al-Ti phase diagram from [2007Wit] is included in the present volume in the evaluation of the Al-C-Ti system.

Solid Phases No ternary compounds exist in the system. Above 882˚C a continuous solid solutions β, (βTi, Ta) is found and above 724˚C a continuous solid solutions ε, (Ti,Ta)Al3 forms [1966Ram, 1983Sri, 1990Abd, 1993Das, 1993Jew]. Single crystals of the (Ti,Ta)Al3 compound were produced by spontaneous crystallization during slow cooling of Al-0.38Ti-0.10Ta (at.%) melt [1990Abd]. The chemical composition of the aluminide crystals separated from melt was Al-75Ti-20Ta (at.%). The crystals had the same crystal structure as ε(h) (TiAl3) and were formed by partial substituting in the lattice Ti atoms by Ta atoms. At 1100˚C the β phase (including the ordered β0 phase) extends up to about 40 at.% Al into the ternary system. In the range of 1350 to 1450˚C it extends to 42–50 at.% Al after [1995Wea, 1993Das, 1993Jew]. The ordering of the bcc phase was established based on the existence of the thermal antiphase boundaries in ternary alloys quenched from 1350 and 1450˚C in [1995Wea]. Such a structure could indicate that these alloys are disordered at high temperature and undergo ordering during cooling. The ordered phase was experimentally revealed at 1100˚C by [1993Jew, 1993Das]. [1993Das] reported that the ordering takes place in the bcc phase (coexisting with other phases) within a composition range of about 40 to 60 at.% Ti and 15 to 60 at.% Al. Near the composition Ti-25Al-25Ta (at.%) the existence of the β0 phase in the ternary region is established. The phase reactions in a Ti-33Al-17Ta (at.%) alloy were examined in detail following solidification and solid-state processing treatments. Differential thermal analysis (DTA) proved that β0-phase in this alloy is stable up to 1205˚C, where it experiences a solid-state order to disorder transformation. The ordering reaction is too fast that it could not be arrested by rapid solidification processing. It is the presence of thermal antiphase boundaries in the microstructure that confirms the solid-state ordering of the β0 phase from the disordered β phase. Taking into account that the β0 phase exists in the binary Al-Ti system at about 1100 to 1430˚C, its range in the ternary system has to adjoin to the Al-Ti side of the Gibbs triangle in the above mentioned temperature interval. Thus, its composition limits depend considerably on temperature. The solubility of the β stabilizer Ta in (αTi) was determined in [1963Luz]. A single-phase region around Ti4Al3Ta, denoted α*, is assumed by [1993Jew] to be a new ternary phase based on α2 (of the close crystal structure). It was observed both in bulk alloys at 1100˚C and also in some of the diffusion couples. There is, however, no confirmation of the existence of this phase by the other authors. Crystallographic data of the Al-Ta-Ti phases and their temperature ranges of stability are listed in Table 2. DOI: 10.1007/978-3-540-88053-0_15 ß Springer 2009

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Invariant Equilibria No experimentally established invariant equilibria in the ternary system are available. A tentative reaction scheme is published by [2005Rag] based on the limited data on monovariant lines on the liquidus, namely β/α, α/γ and γ/η [1992McC, 1991McC1, 1991McC2], the data for the isothermal sections at 1440˚C to 1000˚C [1993Jew, 1995Wea, 1991Boe, 2000Kai] and the boundary binary phase diagrams accepted from [Mas2]. A partial reaction scheme is presented in Fig. 2 for liquid/solid equilibria based only on the accepted tentative liquidus surface projection shown in Fig. 3. Since no data on the temperatures of invariant equilibria in the ternary system exist, in the reaction scheme they were approximated by [2005Rag] based on the binary systems and on the conclusion of [1992McC, 1991McC1, 1991McC2] concerning the direction of the liquidus surface decrease in the ternary system. All the five invariant reactions of Fig. 2 are shown in Fig. 3 without providing temperatures.

Liquidus, Solidus and Solvus Surfaces Limited information on the liquidus surface, obtained by [1991McC1, 1991McC2, 1992McC], concerns only the monovariant l Ð β + α, l Ð α + γ lines and partly l Ð γ + ε. A steep slope of the liquidus β- and α primary crystallization surfaces towards the Al-Ti binary side was concluded. This was concluded from microstructure observations during solidification of ternary alloys in the vicinity of the 50 at.% Al isoconcentrate above 20 at.% Ti. Below 20 at.% Ti an estimated monovariant line l Ð β + σ has been assumed, however not supported by experimental points. No data are available on the solidification of ternary alloys in other parts of the system. From the above mentioned information together with the data on the phase equilibria at 1440˚C by [1993Jew], the existence of two four-phase peritectic reactions was proposed by [2005Rag] who published a schematic liquidus projection and the above tentative reaction scheme including solid state equilibria. The tentative liquidus projection taken from [2005Rag] is presented here in Fig. 3 with amendments made according to the accepted binary diagrams. There is no explicit experimental work describing the equilibria on the solidus surface. In the present work it is assumed that it must be similar to the highest isothermal sections studied in literature by [1995Wea] at 1450˚C and by [1993Jew, 1993Das] at 1440˚C.

Isothermal Sections The isothermal sections at 1450, 1440, 1350, 1330, 1300, 1200, 1100 and 1000˚C were studied experimentally [2000Kai, 1995Wea, 1993Das, 1993Jew, 1991Boe, 1983Sri, 1966Ram]. The experimental results of the above works are in satisfactory agreement with each other. [2005Das] discussed and presented isothermal sections at 1450, 1440, 1350, 1100 and 1000˚C which he derived from the works of [1995Wea, 1993Jew, 1993Das, 1983Sri, 1966Ram]. [2005Rag] considered the above sections at 1440, 1350 and 1100˚C, an isothermal section at 1330˚C constructed from the data of [1991Boe] and partial sections at 1300 and 1200˚C after [2000Kai]. Both review papers, [2005Das] and [2005Rag], apply corrections to Landolt‐Bo¨rnstein New Series IV/11E1

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add missing phase fields or to ensure consistency with the accepted binary diagrams and the phase relations proposed are not in conflict. In this work the isothermal sections given in Figs. 4 to 11 are drawn from the experimental data. They take into account the here accepted edge binary diagrams and the acceptable corrections made in [2005Das] and [2005Rag]. The phase diagrams at 1450 and 1440˚C, Figs. 4 and 5, are constructed after [1995Wea] and [1993Jew, 1993Das]. Both show the three-phase fields σ + β + α, σ + α + δ, α + δ + η and α + η + L and, in addition, at 1440˚C the σ + TaAl + δ field appears because of the existence of the equiatomic tantalum aluminide at this temperature. The solid solution of Al in the α phase reaches up to 85–87 at.% Al at 1450, 1440˚C and then sharply decreases with decreasing temperature after [1995Wea] and [1993Jew, 1993Das], as one can see at 1350 and 1100˚C, Figs. 6 and Fig. 10 respectively. The γ solid solution region at 1450˚ should reach the Al-Ti side unlike the interpretations given by [1995Wea] in his original figure and by [2005Das] in his review. The γ phase is shown to permeate into the ternary system up to about 80 at.% Ta at 1450, 1440 and 1350˚C. Its extension varies only slightly with temperature as one can see from comparison with the data at 1100˚C. The ε phase extension at 1440˚C given by [1993Das, 1993Jew] had to be corrected to agree with that at 1450 and 1350˚C by the same authors. At 1440˚C and below, equilibria with the participation of tantalum monoaluminide had to be added. Below 1345˚C the Ta5Al7 phase exists in the isothermal section and below 1226˚C the Ta2Al3 phase appears. But the Al69Ta39 phase (δ) disappears at 1183˚C from the phase equilibria because of the eutectoid decomposition of this phase. For the equilibria involving α and δ a principal change takes place in the temperature interval of 1440 to 1330˚C. The α + δ equilibria change for σ + ε and σ + γ and these equilibria are retained down to the lowest temperature under investigation (1000˚C [1966Ram]). The existence of the γ + σ equilibrium and a significant temperature dependence of the tantalum solubility in γ were confirmed by [1992McC] who observed the decomposition of γ into a two-phase (γ + σ) lamellar structure. A partial isothermal section at 1330˚C after [1991Boe] is given in Fig. 7. At 1350˚C the ε phase reaches up to the Al-Ti binary. At the high Ti border it is in equilibrium with γ and at the high Al border with L, contrary to [1995Wea], who assumed that L is in equilibrium with γ. Other equilibria missed out by 1995Wea] concern the equilibria with the TaAl and δ phases are shown in the isothermal section at 1350˚C (Fig. 6). That means the two three-phase fields σ + δ + ε and σ + δ + TaAl have been added tentatively by dashed lines. Also the β0, the ordered bcc phase, which exists at this temperature in the binary Al-Ti is added to Fig. 6. The β0 ternary solid solution was found by [1993Das, 1993Jew] at 1100˚C, just below the temperature where it disappears in the binary Al-Ti system. The 1100˚C isothermal section in Fig. 10 is based on the experimental data of [1993Jew, 1993Das] with the following corrections: the position of the β (β0)-phase fields is adjusted to match the accepted Al-Ti diagram; equilibria with the η and ζ phases at the Al-Ti side have been added as well as the equilibria with the Ta5Al7 and Ta2Al3 phases near the Al-Ta side. The main changes in the phase relationships at 1100˚C - with respect to higher temperatures - are (i) that the α phase field became very small (supported by [1983Sri]), (ii) an extended α2 field exists and (iii) an equilibrium exists between γ and β0 (the γ + β0 field in Fig. 10) instead of α + σ at higher temperatures. If one compares the isothermal sections at 1300 and 1200˚C, Figs. 8 and 9, published by [2000Kai] a transformation α + σ Ð γ + β0 must have taken place in this temperature interval. DOI: 10.1007/978-3-540-88053-0_15 ß Springer 2009

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The isothermal section at 1000˚C given in Fig. 11 is based on the preliminary data published by [1966Ram], again amended to match the accepted binary diagrams. A remarkable difference of the phase relationships from that at 1100˚C is the existence of the σ + α2 equilibrium instead of β0 + γ and the absence of the ordered bcc (β0) phase, which is in agreement with its absence in the binary Al-Ti. The extension of the γ filed into the ternary system seems to be too large by [1966Ram] and has been reduced slightly in Fig. 11 to nearly the same extension as at high temperatures. The phase equilibria at 1100˚C, reported by [1983Sri] and later accepted by [1993Kub], look like those at 1000˚C published by [1966Ram]. They are significantly different from those established by [1993Jew, 1993Das]. The reason of this contradiction is not quite clear, although it is suspected that contamination during the sample preparation could be the reason. The initial materials were of about the same purity and the methods of preparation and investigations were almost the same. The alloys in [1983Sri] were prepared by sintering the powder mixtures of the components as well as by arc melting and subsequently annealed in quartz ampoules. The duration of thermal treatment was at least 170 h. As stated in [1983Sri], the alloys could be contaminated by silicon during the preparation, because the impurity phases Ti5Si3 and Ta5Si3 were observed. In contrast to [1983Sri], [1966Ram] prepared specimens by arc-melting and heat treated them for 7 d in evacuated quartz containers, encapsulated by Ta-foil with subsequent quenching. Since the results of [1983Sri] do not follow the general trend with changing temperature, these data are not accepted in the present evaluation.

Notes on Materials Properties and Applications The works devoted to the study of various mechanical, structural and physical properties [1993Has, 1995Muk, 1999Hao, 2001Gar, 2002Red, 2005Leg, 2006Joh, 2006How] are summarized in Table 3. Electrical resistance measurements and hardness tests as well as X-ray and metallographic analyses were carried out by [1963Luz] to determine the solubility of the β stabilizer Ta in (αTi). Mechanical properties of Ti-46.4Al-2.5Ta (at.%) alloy were investigated in [1993Has]. With increasing temperature from 800 to 1100˚C the elongation of sample increased from 16 to 104%. At the same time yield stress decreased from 394 to 28.4 MPa. The X-ray technique was used together with micro-indentation to measure average residual stresses and local hardness of the alloy of Ti-47Al-2Ta (at.%) composition [2001Gar]. The measured values of residual stresses were compared to a composite cylinder model, which incorporates the effect of creep. In general, the residual stresses decrease as heat treatment temperature increases, reaching a minimum and then increases as the temperature and/or time of treatment increases. The fiber-matrix interface shows no degradation during the heat treating process. Such results compare well with the composite cylinder model once the effect of creep is incorporated. The analysis of the micro-hardness results indicate an increase in hardness in regions close to the fiber-matrix interface at 593˚C but a softening occurs at 815˚C. This behavior is reversed once the heat treatment temperature is increased to 982˚C. At this temperature, the hardness at the interface increases for prolonged heat treatment duration. Landolt‐Bo¨rnstein New Series IV/11E1

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Miscellaneous Mechanical alloying of blended elemental Ti-25Al-25Ta (at.%) was investigated in [1995Muk]. Mechanical alloying led initially to the formation of a disordered bcc phase with a = 332 pm, and at still longer times a fully amorphous phase. The disordered bcc phase transformed to the B2 (the ordered bcc) structure after heat treatment at 600˚C and to an orthorhombic phase (with a = 608.9 pm, b = 857.0 pm, c = 498.0 pm) after a heat treatment at 800˚C. The site occupancies of Ta in γαffl (TiAl) alloys with different compositions, and in Ti3Al with the compositions of Ti - 26 at.% Al - (1 to 2) at.% Ta, were measured by the atom location channeling enhanced microanalysis (ALCHEMI) method [1999Hao]. For TiAl alloys, the results show that Ta atoms invariably occupy Ti sites. For Ti3Al alloys Ta atoms occupy Ti sites too. The experimental results were interpreted in terms of a Bragg-Williams-type model and bond-order data were obtained from electronic structure calculation. Qualitative agreement between the model and measurements was reached. A theoretical model that relates the (γ + α2) two-phase equilibrium in ternary Al-Ta-Ti alloys to the substitution behavior of alloying elements in the two ordered phases has been suggested in [1999Yan]. The method used experimentally measured occupation probabilities of the alloying elements on different sublattices as input, and allowed estimates of the mole fractions of the two phases for a given alloy composition. The review summarized theoretical and experimental investigations of sublattice substitution of alloying elements in γαffl phase (TiAl) presented in [2000Yan]. The isothermal oxidation behavior of a ternary Ti-25Al-18Ta (at.%) Ta intermetallic alloy has been investigated in pure oxygen over the temperature range of 850 to 1100˚C [2002Red]. The oxidation kinetics was found to follow a parabolic rate. Effective activation energy of 259 kJ·mol–1 was deduced from the oxidation data. The oxidation products were a mixture of TiO2, the main component, Al2O3 (alumina), and small amounts of tantalum oxide. The addition of Ta to Ti3Al alloy decreased the oxidation rate of the alloy. However, the oxidation scale was not compact and exhibited significant spallation especially at high temperatures. The microchemistry and morphology of the oxide layer produced on a single γαffl phase (TiAl) alloy (Ti-52.1Al-2Ta (at.%)) following anneals at 550 and 900˚C in low-pressure oxygen and hydrogen environments were studied in [2005Leg]. The oxide structure on the electropolished surface consisted of two layers, an outer one of Al2O3 and an inner one of TiO2. Annealing in low partial pressures of oxygen retained the same oxide structure but increased the total thickness. After the low-pressure oxygen anneals at 550˚C, the oxide surface was smooth, whereas after anneals at 800˚C and above the surface consisted of small, round particles, whose size and density increased with increasing annealing temperature and oxygen partial pressure. In contrast, the oxide structure produced by annealing in a hydrogen environment after the low-pressure oxygen treatment consisted of an outer layer of TiO2 and a sub-oxide of Al2O3. The morphology of this oxide consisted of elongated, rod-like particles, whose size increased with annealing temperature. Microstructural selection during the directional solidification of binary γαffl (TiAl) alloys grown from Ti-52Al-8Ta (at.%) seeds was examined [2006Joh]. By using a seed crystal, the high-temperature hcp α phase can be correctly oriented so that an aligned lamellar γαffl (TiAl)/ Ti3Al microstructure results from the subsequent solid-state transformations upon cooling. From the equilibrium phase diagram, primary bcc β phase solidification is expected in the compositional range for which the binary γ (TiAl) alloys were successfully seeded. Thus successful crystal growth of γ (TiAl) was found to be dependent upon the undercooling DOI: 10.1007/978-3-540-88053-0_15 ß Springer 2009

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necessary for β nucleation. From these data, a microstructural selection map for the seeded growth of the α phase was constructed. The atomic-level structural, compositional, kinetic and mechanistic aspects of phase transformations that occur during growth of product phases in fcc to hcp reactions were investigated in [2006How] for Ti-48Al- 2Ta (at.%) alloy.

. Table 1 Investigations of the Al-Ta-Ti Phase Relations and Structures Temperature/Composition/Phase Range Studied

Reference


[1963Luz]

Electrical resistance measurements, hardness tests, X-ray diffraction (XRD), metallographic analysis

solubility of Ta in (αTi).

[1966Ram]

XRD

isothermal section at 1000˚C

[1983Sri]

XRD


[1990Abd]

XRD, scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS)

crystal structure of ε phase

[1991Boe]

metallographic analysis, transmission alloy Al-25Ti-25Ta (at.%) at 1200 to electron microscopy (TEM), SEM, electron 1550˚C, isothermal section at 1330˚C probe microanalysis (EPMA), differential thermal analysis (DTA)

[1991Das]

TEM, EPMA, DTA, diffusion couple technique


[1991McC1] SEM, TEM, EDS

undercooled and splat quenched alloy Ti-15Ta-48Al (at.%), solidification path

[1991McC2] XRD, SEM, TEM, EDS

alloys at 50 at.% Al, isothermal section at 1100˚C, partial liquidus surface

[1991Wea]

metallographic analysis, XRD, TEM, DTA

alloys at 26 to 70 at.% Al, 8 to 54 at.% Ti, 5 to 50 at.% Ta, isothermal section at 1450˚C, morphology

[1992Kim]

metallographic analysis, XRD, EPMA

alloy Ti-46.4Al-2.5Ta (at.%) at 1000 to 1300˚C, phase constitution, morphology

[1992McC]

XRD, SEM, TEM, EDS

alloys Ti-0 to 27 at.% Ta-45 to 70 at.% Al, partial liquidus surface, morphology and phase constitution at 1200˚C

[1992Per]

SEM, EPMA, DTA, XRD


[1992Wea]

metallographic analysis, SEM, EPMA, TEM, alloys at 64 to 71 at.% Al, 7 to 13 at.% Ti, XRD 16 to 32 at.% Ta, isothermal section at 1450˚C

[1993Das]

TEM, SEM, EPMA, DTA, XRD study of diffusion couple and bulk alloy samples


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[1993Jew]

metallographic analysis, TEM, SEM, EPMA, isothermal sections at 1100 and 1440˚C DTA, XRD

[1993Has]

metallographic analysis, XRD

[1995Wea]

metallographic analysis, SEM, EPMA, TEM alloys at 38 to 55 at.% Al, 19 to 30 at.% Ti, 20 to 38 at.% Ta, solid state phase transformations, isothermal sections at 1350 and 1450˚C, morphology

[2000Kai]

SEM, EPMA

alloy Ti-47Al-6Ta (at.%) at 1000 to 1300˚C, isothermal sections at 1000, 1200 and 1300˚C

[2003Das]

DTA, TEM, splat-quenching technique

ordering of the bcc-phase

alloy Ti-46.4Al-2.5Ta (at.%) at 1100˚C, phase constitution, morphology



(Al) < 660.5

cF4 Fm3m Cu

β, (βTi,Ta) Ti1–x–yTaxAly < 3020

cI2 Im3m W

(Ta) < 3020 (βTi) 1670 - 882 β0, Ti1–x–yTaxAly cP2 < 1427 Pm3m CsCl

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Lattice Parameters [pm] a = 404.96 a = 405.2

Comments/References pure Al at 25˚C [Mas2] Solubility limits: at 0% Ti, 0.04 at.% Ta [1972Fer], at 0% Ta, 0.8 at.% Ti [2007Wit] or 0.6 at.% Ti [1992Kat]

a = 330.30

complete solid solution [Mas2], 0 < x < 1 at T = 1440˚C, x = 0 to 1 and y = 0 to 0.5 [1993Das]; at T = 1350˚C, x = 0 to 1 and y = 0 to 0.42 [1995Wea]; at T = 1100˚C, x = 0 to 1 and y = 0 to 0.37 [1993Das]; pure Ta at 25˚C [Mas2]

a = 330.65

pure Ti at 900˚C [Mas2]

-

at T = 1100˚C, y = 0.21 to 0.37 and x = 0 to 0.45 [1993Das] at x = 0, y = 0.22 to 0.42 1427 < T < 1114 [2007Wit]

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. Table 2 (continued) Phase/ Temperature Range [˚C] α, (αTi) Ti1–x–yTaxAly 1491 - 1119 and < 1159



hP2 P63/mmc Mg

a = 295.08 c = 468.35

α2, Ti3Al Ti1–x–yTaxAly < 1189

hP8 P63/mmc Ni3Sn

α*, Ti1–x–yTaxAly crystal structure close to α2

γ, TiAl Ti1–x–yTaxAly < 1463

ζ, Ti2Al5 Ti1–x–yTaxAly 1432 - 976

1d-APD

tP28 P4/mmm “Ti2Al5”


a = 580.6 c = 465.5 a = 574.6 c = 462.4

at x = 0: 17 to 38.5 at.% Al [2007Wit]; at 22 at.% Al [2004Sch] at 38 at.% Al [2004Sch] at T = 1100˚C, y = 0.33 to 0.48 and x = 0.07 to 0.09 [1993Das] Ternary phase? The existence needs confirmation

-

a = 400.0 c = 407.5 a = 398.4 c = 406.0 set of tetragonal superstructures* based on AuCu-lattice;

at T = 1440˚C, x = 0.28 and y = 0.60 [1993Das]; at T = 1350˚C, x = 0.22 and y = 0.51 [1995Wea]; at x = 0.02 and T = 1100˚C, y = 0.2 [1993Das]; pure Ti, at 25˚C [V-C2] at x = 0, y = 0 to 0.28 [2007Wit]; at x = 0 and T = 1189˚C, y = 0.32 [2007Wit]; at x = 0 and T = 1456˚C, y = 0.5 [2007Wit]; 0.27 < y < 0.48 and 0 < x < 0.09 at 1100˚C (including α2*)

tP4 P4/mmm AuCu

TiAl

Comments/References

at T = 1100˚C, 0<x 99.85 mass% Nb and amorphous > 97 mass% B) and B4C which was annealed in vacuum at 2000˚C to reduce C content to 0.2 mass% free C. Specimens in form of cylinders were compacted with aid of 12% aqueous starch solution, presintered at 2100˚C for 2 h in vacuum prior. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

Investigation of the quasieutectic system NbB2+‘B4C’ with eutectic point at 35–37 mol% NbB2 and TE of 2250˚C on 12 samples.


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[2001Hil]

Compounds forming a structural series Determination of the crystal structures (NbB)2(NbB2)n(NbC)m were synthesised (Nb3B3C, Nb4B3C2, Nb7B6C3 and Nb7B4C4) from Cu/Al flux in Al2O3 crucibles under from X-ray counter data. argon starting from Cu:Al:B:C:Nb ratios 20:5:2:1:4 (Nb3B3C); 20:5:0.2:0.1:0.2 (Nb4B3C2); 20:5:2:1:2 (Nb7B6C3); 20:5:0.1:0.1:0.2 (Nb7B4C4). The mixtures were heated to 1600˚C at 400˚/h, kept for 24 h and cooled to 1500˚C at 100˚/h, to 1100˚C at 1˚/h and to RT at 150˚/h. Single crystals were obtained after dissolution of the flux in halfconcentrated HNO3. From EDX only Nb was found as a metal. In all cases NbC was found as a side product.

[2004Kor]

More than 30 ternary samples prepared with various techniques (A, B, C) from high purity powders of components: niobium (99.9 and 99.85 mass%), boron (Ventron GmbH, D), boron carbide B4C (Johnson Matthey & Co, UK) and carbon (with purity of 99.99 mass%). For melting point measurements by the Pirani technique [1923Pir] specimens were prepared in form of cylindrical polycrystalline rods by sintering at 2000˚C of ceramic green bodies from isostatically compressed blends of Nb, B4C, B and C powders. Powder blends were compacted in form of the final Pirani shape in steel dies using a small amount of xylol as densification aid. The green bodies (with a lateral hole directly pressed into the sample) were slowly heated (to prevent violent self-heating via exothermic reactions) in high vacuum to 1600˚C for 12 to 15 h. For LOM and EMPA, several alloys were prepared by argon arc melting of sintered pellets

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

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Melting points were determined on bout 4 Pirani samples of each composition in a precise micro-pyrometer calibrated for the temperature region from 2000 to 2900˚C. The microstructure of the cast alloys was inspected using light optical microscopy (LOM) on flat surfaces prepared by grinding (SiC-paper) and polishing the resin-mounted alloys using diamond pastes down to 1/4 μm grain size. Quantitative composition analyses were performed on a CAMEBAX SX50 wavelength dispersive spectrograph (EMPA) comparing the X-ray emissions of the three elements in the alloys with those from elemental standards of Mo, Si and LaB5.85 for boron after applying a deconvolution and ZAF-correction procedure. Determination of liquidus and solidus in the entire range of compositions. Determination of isothermal section at 1750˚C. Determination of reaction isotherms for reactions involving liquid. Derivation of Schulz-Scheil diagram for the entire diagram. Revision of isopleths NbB2-NbC1–x, ‘B4C’(B4.5C)-NbB2


B–C–Nb

20




[2004Pad] Rods of 8–10mm diameter and 50 to 100 mm length were prepared from powder compacts (99.5 mass% NbB2 and 99.9% B4.3C) annealed at 1500˚C in vacuum prior to zone melting in an induction furnace under 0.2 MPa argon. Compositions chosen were: Nb11B76C13 (46 mol% NbB2) and Nb7.5B77C15.5 (33 mol% NbB2). XPD, SEM, EMPA.

Directional crystallization of B4C-NbB2 eutectic compositions investigated by XPD, SEM and EMPA revealing NbB2 needles (whiskers of 1–2 mm diameter)

[2005Tan] Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)) the liquidus and solidus curves to the L + (βB) field has been derived via chemical analysis

Confirmation of peritectic type of reaction L + B4+xC Ð (βB) via determination of the liquidus and solidus curves to the L + (βB) field





Comments/References

(Nb) < 2469

cI2 Im 3m W

a = 330.04

at 25˚C [Mas2]

(βB) < 2092

hR333 R3m βB

a = 1093.30 c = 2382.52

[1993Wer]

a = 1093.02 c = 2381.66 a = 1092.2 c = 2381.4 a = 1091.91 c = 2382.24

pure B [1976Lun]

(αB)

hR36 R3m αB


a = 490.8 c = 1256.7

MSIT1

at 1.1 at.% C [1993Wer], linear da/dx, dc/dx at NbB99.5 [1992Rog] presumably metastable phase, preparation below 1000˚C [1971Amb] pure B, single crystal [1994Cha]

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20

B–C–Nb





(C)gr hP4 a = 246.12 < 3827 P63/mmc c = 670.90 (sublimation point) C (graphite) a = 246.023 c = 671.163 a = 246.75 c = 669.78

Comments/References at 25˚C [Mas2]

[1967Low] at 2.35 at.% Cmax (2350˚C), linear da/ dx, dc/dx [1967Low]

(C)d

cF8 a = 356.69 Fd 3m C (diamond)

at 25˚C, 60 GPa [Mas2]

‘B4C’ < 2450

hR45 R 3m B13C2

a = 565.1 to 560.7 c = 1219.6 to 1209.5

9 to 20 at.% C [1990Ase]

a = 556.0 c = 1212.0 a = 556.1 c = 1212.0

sample quenched from 2400˚C [1987Ord] sample ‘B4C’ + 14.5 mol% NbB2 quenched from 2420˚C [1987Ord]

B25C

tP68 P 42m or P42/nnm B25C

a = 875.3 ± 0.4 c = 509.3 ± 1.5

[V-C2] also B51C1, B49C3; all metastable? defect structure tP52 [1998Rog]

Nb3B2 < 2080

tP10 P4/mbm U3Si2

a = 619.79 c = 329.26

[1992Rog] oxygen stabilized? [1985Zak]

NbB < 2917

oC8 Cmcm CrB

a = 329.74 b = 872.38 c = 316.69 a = 330.2 b = 875.5 c = 316.8 a = 330.0 b = 873.3 c = 316.6 a = 329.61 b = 872.24 c = 316.53

[1992Rog]

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

MSIT1

as-cast, at 32.9 at.% B [1985Zak] at 32.9 at.% B, 25 h at 1950˚C [1985Zak] [1991Oka]


B–C–Nb

20


Phase/ Temperature Range [˚C] Nb5B6 < 2870

Nb3B4 < 2935

Pearson Symbol/ Space Group/ Prototype oC22 Cmmm V5B6

oI14 Immm Ta3B4

Nb2B3

oC20 Cmcm V2B3

NbB2 < 3036

hP3 P6/mmm AlB2

Lattice Parameters [pm] a = 315.30 b = 2227.44 c = 330.49 a = 315.67 b = 2276.7 c = 330.34 a = 314.1 b = 2275.6 c = 330.6 a = 314.51 b = 1410.62 c = 330.19 a = 314.28 b = 1407.6 c = 330.33 a = 330.58 b = 1948.1 c = 312.93 a = 311.26 c = 326.27 a = 308.61 c = 330.69 a = 311.15 c = 326.57 a = 310.37 c = 332.37 a = 311.2 c = 327.0 a = 312.3 c = 329.0 a = 314.0 c = 331.5 a = 309.5 c = 330.0 a = 308.9 c = 330.2 a = 310.0 c = 330.0


MSIT1

Comments/References [1992Rog]

[1991Oka]

as-cast, together with Nb3B4, τ1 and NbC1–x [2004Kor] [1992Rog]

[1991Oka]

[1991Oka]

65 to 70 at.% B Nb rich [1992Rog] B rich [1992Rog] Nb rich [1991Oka] B rich [1991Oka] at 227˚C [V-C2] at 727˚C [V-C2] at 1227˚C [V-C2] for NbB2, quenched from 3000˚C [1987Ord] for samples NbB2 and NbB2 + 1 mol % NbC1–x quenched from 2600˚C [1977Ord] for sample NbB2 + 7.7 mol% ‘B4C’, quenched from 2620˚C [1987Ord]

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B–C–Nb


Phase/ Temperature Range [˚C] Nb2C(h2) 3080 - 2450



hP4 P63/mmc defect NiAs

28 to 35.5 at.% C; labeled as γNb2C [1996Len, 1997Len, 1998Wie] defect structure hP3 [1998Rog] for NbC0.36 quenched from 1750˚C [1963Rud] for Nb2C quenched from 1750˚C [1963Rud] at 11.11 at.% C, as-cast [1985Zak]

a = 311.6 c = 495.8 a = 312.4 c = 496.8 a = 311.9 c = 495.6 a = 311.5 c = 495.4 Nb2C(h1) 2530 - 1195

hP9 P31m W2C1)

Nb2C(r) < 1195

oP12 Pnma Nb2C42) (MnO2)

Nb4C3–x < 1575 ± 25

hR24 R 3m V4C3

Nb6C5 < 1050

mC22 C2/m Nb6C53)

hP36 or P31 Nb6C5

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

Comments/References

at 11.11 at.% C, 150 h annealing at 1200˚C [1985Zak]

a = 541.69 c = 479.19

27 to 36.6 at.% C; labeled as βNb2C [1996Len, 1997Len, 1998Wie] [V-C2]

a = 1091.0 b = 309.54 c = 497.46

at 34.6 to 36.6 at.% C; labeled as αNb2C [1967Rud] (impurity stabilized?)

a = 314 ± 1 c = 3010 ± 10

40.1 to 40.7 at.% C [1996Len, 1997Len, 1998Wie] [V-C2] defect structure tP52 [1998Rog]

a = 544.7 b = 943.5 c = 544.7 β = 109.47˚ a = 546.4 c = 1542.2

MSIT1

at 45 at.% C [Mas2] [V-C2]

[V-C2]


B–C–Nb

20


Phase/ Temperature Range [˚C] NbC1–x < 3600

Pearson Symbol/ Space Group/ Prototype cF8 Fm 3m NaCl

Lattice Parameters [pm] a = 443.0 a = 447.0 a = 446.81 a = 448.44 a = 443.13 a = 446.90 to 447.36

* τ1, Nb3B3C < 2970

oC28 Cmcm Nb3B3C

a = 326.47 b = 2871.0 c = 312.85 a = 325.9 b = 2875.5 c = 312.7

Comments/References 40 to 49.6 at.% C [Mas2] at 41 at.% C, quenched from 1750˚C [1963Rud] at 48 at.% C, quenched from 1750˚C [1963Rud] 48 at.% C, T = 20˚C [V-C2] 48 at.% C, T = 616˚C [V-C2] 41 at.% C, T = 20˚C [V-C2] from samples NbC1–x + NbB2, quenched from 2600˚C; linear increase from 0 to 7 mol% NbB2 [1977Ord] [2001Hil]

as-cast [2004Kor]

* τ2, Nb4B3C2

oC36 Cmcm Nb4B3C2

a = 322.88 b = 3754.4 c = 313.32

[2001Hil] metastable?

* τ3, Nb7B6C3

oC32 Cmmm Nb7B6C3

a = 313.41 b = 3316.1 c = 324.28


* τ4, Nb7B4C4

oI30 Immm Nb7B4C4

a = 315.442 b = 321.66 c = 3226.0


partially ordered V2N type, cited by various authors as partially ordered εFe2N type [1963Rud]. [1967Rud] indexed Nb2C (r) on the basis of the ζFe2N type with cell parameters a = 1089.5, b = 496.8, c = 1235 pm. This phase was assumed to be impurity (oxygen?) stabilized and was thus removed from the diagram in [Mas2]. 3) Nb6C5 as claimed by [1991Gus], was said to be a NaCl-type derivative superlattice structure with symmetry C2, C2/m or P31. 1) 2)


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B–C–Nb


T [˚C]

Type

Phase

B

C

Nb

L Ð Nb3B4 + τ1

2940

e2 (max)

L

44.5

12.5

43

L Ð NbB2 + τ1

2910

e3 (max)

L

44.4

13.2

42.4

L Ð NbB2 + τ1 + Nb3B4

2900 ± 15

E1

L

44.8

12.5

42.7

L Ð NbB2 + NbC1–x

2888 ± 19

e4 (max)

L

43

16

41

L Ð Nb3B4 + NbC1–x

2830 ± 30

e6 (max)

L

44.

10

46

L Ð NbB + NbC1–x

2800

e7 (max)

L

40

8.5

51.5

NbB2

66

1

33

NbC1-x

4

40

56

14.5

43

L Ð τ1 + NbC1–x

2790 ± 10

e8 (max)

L

42.5

L+Nb3B4 Ð Nb5B6 + NbC1–x

2772 ± 13

U1

L

44

9.5

46.5

L Ð Nb5B6 + NbC1–x + NbB

2728 ± 12

E2

L

42.5

8.0

49.5

L Ð NbB2 + (C)gr

2707 ± 9

e9 (max)

L

47

29

24

L Ð NbB2 + τ1 + NbC1–x

2700 ± 20

E3

L

43

15.5

41.5

L Ð NbB2 + NbC1–x + (C)gr

2570 ± 10

E4

L

40

31.5

28.5

L + NbC1–x Ð NbB + Nb2C

2375 ± 15

U2

L

28.5

14.0

57.5

L Ð NbB2 + ‘B4C’

2290 ± 25

e12 (max)

L

78

13

9

L Ð Nb3B4 + τ1 + NbC1–x

2340 ± 20

E5

L

43

12.5

44.5

L Ð NbB2 + ‘B4C’ + (C)gr

2243 ± 11

E6

L

62.5

22.5

15

L + ‘B4C’ Ð NbB2 + (βB)

< 2103

U3

L

98

1

1

L Ð (Nb) + NbB + Nb2C

2060

E7

L

18

7

75

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

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B–C–Nb

20

. Fig. 1 B-C-Nb. The C-Nb phase diagram


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B–C–Nb

. Fig. 2 B-C-Nb. The B-Nb phase diagram. Phase boundaries from [1985Zak] are shown as dash lines, from [1963Rud] as dotted lines

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

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B–C–Nb

20

. Fig. 3 B-C-Nb. The NbB2-NbC1–x phase diagram. Dashed lines from [1977Ord]; dotted lines from [2004Kor]


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B–C–Nb

. Fig. 4 B-C-Nb. The ‘B4C’ (B4.5C)-NbB2 phase diagram

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

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20

. Fig. 5a B-C-Nb. Reaction scheme, part 1

B–C–Nb


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B–C–Nb

. Fig. 5b B-C-Nb. Reaction scheme, part 2

20

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B–C–Nb

20

. Fig. 6 B-C-Nb. Liquidus surface projection


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22

20

B–C–Nb

. Fig. 7 B-C-Nb. Isothermal section at 1750˚C

DOI: 10.1007/978-3-540-88053-0_20 ß Springer 2009

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B–C–Nb

20

References [1923Pir] [1952Gla] [1955Bre] [1963Rud]

[1965Lev]

[1966Rud]

[1967Low] [1967Rud] [1971Amb] [1975Kry]

[1976Lun] [1977Ord]

[1978Ord]

[1980Ord]

[1983Sch]

[1984Hol]

[1984Ord]

[1985Zak]

[1987Ord]

Pirani, M., Alterthum, H., “Method for the Determination of the Melting Point of Metals which Fuse at High Temperatures” (in German), Z. Elektrochem., 29, 5–8 (1923) (Phase Relations, Experimental, 5) Glaser, F.W., “Contribution to the Metal-Carbon-Boron Systems”, Trans. AIME - J. Metals, (4), 391–396 (1952) (Crys. Structure, Experimental, Electr. Prop., 19) Brewer, L, Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–405 (1955) (Phase Relations, Thermodyn., Experimental, Review, 19) Rudy, E., Benesovsky, F., Toth, L., “Studies of the Ternary Systems of the Group Va and VIa Metals with Boron and Carbon” (in German), Z. Metallkd., 54, 345–353 (1963) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 43) Levinskii, Yu.V., Salibekov, S.E., Levinskaya, M.K., “Interaction of Diborides of V, Nb, Ta with Carbon” (in Russian), Poroshk. Metall., (Kiev), (5), 66–69 (1965) (Phase Diagram, Phase Relations, Experimental, *, 5) Rudy, E., Windisch, S., “Ternary Phase Equilibria in Transition Metal-Boron-Carbon-Silicon Systems”, Tech. Rep. AFML-TR-65–2, Air Force Materials Laboratory, Wright-Patterson Air Force Base OH, Part I, Vol. X, 1–103 (1966) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, *, 64) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Rudy, E., Brukl, C.E., “Lower-Temperature Modifications of Nb2C and V2C”, J. Am. Ceram. Soc., 50, 265–268 (1967) (Crys. Structure, Experimental, 14) Amberger, E., Ploog, G., “Formation of Pure Boron Lattice” (in German), J. Less-Common Met., 23, 21–31 (1971) (Crys. Structure, Experimental, 18) Krylov, Y.I., Bronnikov, V.A., Krysina, V.G., Pristavko, V.V., “On the Possibility of Creating Thermite Mixtures Based on Compositions of Boron Carbide (B4C) Materials and Silicon Carbide Materials” (in Russian), Poroshk. Metall., (Kiev), (12), 57–60 (1975) (Morphology, Experimental, Phys. Prop., 5) Lundstro¨m, T., Tergenius, L.E., “On the Solid Solution of Copper in β-Rhombohedral Boron”, J. LessCommon Met., 47, 23–28 (1976) (Crys. Structure, Experimental, 10) Ordan’yan, S.S., Stepanenko, E.K., Unrod, V.I., “Reactions in the System NbC-NbB2”, Inorg. Mater., 13(2), 312–314 (1977), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 13(2), 373–375 (1977) (Crys Structure, Phase Diagram, Experimental, Mechan. Prop., Phase Relations, 4) Ordan’yan, S.S., Stepanenko, E.K., Unrod, V.I., “Reactions in the NbC-NbB2 System” (in Russian), in “Vysokotemp. Boridy Silitsidy”, Kosolapova, T.Ya. (Ed.), Naukova Dumka, Kiev, USSR, 64–67 (1978) (Crys Structure, Phase Diagram, Experimental, Mechan. Prop., Phase Relations, 4) Ordanyan, S.S., “Laws of Interaction in the Systems MIV,VC - MIV,VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411 (1980) (Thermodyn., Experimental, Theory, Electronic Structure, 14) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Phase Diagram, Phase Relations, Review, Mechan. Prop., 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Ordan’yan, S.S., Stepanenko, E.K., Sokolov, N.V., “Strength of Sintered Niobium Carbide - Niobium Boride (NbC-NbB2) Composite Materials” (in Russian), Izv. Vyss. Uchebn. Zaved., Khim., Khim. Tekhnol. USSR, 27(10), 1201–1203 (1984) (Morphology, Experimental, Mechan. Prop., 4) Zakharov, A.M, Pshokin, V.P, Ivanova, E.I., “Niobium Corner of the System Nb-B-C”, Russ. Metall., (5), 192–195 (1985), translated from Izv. Akad. Nauk SSSR, Met., (5), 193–196 (1985) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, *, 10) Ordan’yan, S.S, Dmitriev, A.I., Bizhev, K.T., Stepanenko, E.K., “Methods of Examination and Properties of Powder Material Interaction in B4C-Me(V)B2 Systems”, Sov. Powder Metall. Met. Ceram., 298(10), 834–836 (1987), translated from Poroshk. Metall. (Kiev), 298(10), 66–69 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., *, 5)


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24

20 [1990Ase]

[1991Gus]

[1991Oka]

[1992Rog]

[1993Ord]

[1993Wer]

[1994Cha]

[1994McH]

[1996Kas] [1996Len]

[1997Len]

[1998Kas]

[1998Rog]

[1998Wie]

[2000Gus] [2001Bur] [2001Hil] [2003Bor]

[2004Kor]

B–C–Nb Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides” Freer, R. (Ed.), Dordrecht: Kluwer Academic Publishers, 97–111 (1990) (Crys. Structure, Experimental, Review, 14) Gusev, A.I. “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State., 32(9), 1595–1599 (1991) (Phase Diagram, Phase Relations, Thermodyn., Theory, *, 18) Okada, S., Hamano, K., Lundstroem, T., Higashi, I., “Crystal Growth of the New Compound Nb2B3, and the Borides NbB, Nb5B6, Nb3B4 and NbB2 Using the Copper-flux Method” in “AIP Conference Proceedings 231 on Boron-rich Solids”, Albequerque, USA, 1990, New York, AIP, 590–593 (1991) (Crys. Structure, Experimental, 12) Rogl, P., “The System B-N-Nb” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J. (Eds.), ASM, Materials Park, Ohio, USA, 68–72 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, *, 6) Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, (1), 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, Electronic Structure, 18) Werheit, H., Kuhlmann, U., Laux, M., Lundstroem, T., “Structural and Electronic Properties of Carbon-doped β-Rhombohedral Boron”, Phys. Stat. Sol. B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, Electronic Structure, 51) Chakrabarti, D.J., Laughlin, D.E., “B-Cu (Boron-Copper)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 74–78 (1994) (Review, Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., 24) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 189–190 (1994) (Phase Diagram, Phase Relations, Review, 2) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Lengauer, W., Wiesenberger, H., Joguet, M., Rafaja, D., Ettmayer, P., “Chemical Diffusion in Transition Metal-Nitrogen Systems” in “The Chemistry of Transition Metal Carbides and Nitrides”, Oyama, S.T. (Ed.), Oxford, Blacky Academic, 91–106 (1996) (Phase Diagram, Phase Relations, Experimental, *, 29) Lengauer, W., Wiesenberger, H., Mayr, W., Bidaud, E., Berger, R., Ettmayer, P., “Phase Stabilities of Transition Metal Carbides and Nitrides Investigated by Reaction Diffusion”, J. Chim. Phys., 94, 1020–1025 (1997) (Morphology, Experimental, Interface Phenomena, *, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Niobium” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, 197–213 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, #, 25) Wiesenberger, H., Lengauer, W., Ettmayer, P., “Reaction Diffusion and Phase Equilibria in the V-C, Nb-C, Ta-C and Ta-N systems”, Acta Mater., 46(2), 651–666 (1998) (Phase Diagram, Experimental, Interface Phenomena, *, 30) Gusev, A.I., “Order-Disorder Phase Transformations in Strongly Non-Stoichiometric Compounds” (in Russian), Physics Usbekhi, 3(1), 1–4, 34–37 (2000) (Phase Diagram, Thermodyn., Theory, *, 12) Burkhanov, G.S., “Structural Materials on the Basis of Rare Materials” (in Russian), Metally, (5), 57–61 (2001) (Morphology, Review, Phys. Prop., 46) Hillebrecht, H., Gebhardt, K., “Crystal Structures from a Building Set: the First Boride Carbides of Niobium” (in German), Angew. Chem., 113(8), 1492–1495 (2001) (Crys. Structure, Experimental, 17) Borges, L.A., Jr., Coelho, G.C., Nunes, C.A., Suzuki, P.A., “New Data on Phase Equilibria in the Nb-rich Region of the Nb-B System”, J. Phase Equilib., 24(2), 140–146 (2003) (Crys. Structure, Experimental, *, 14) Korniyenko, K., Rogl, P., Velikanova, T., Leithe-Jasper, A., Bohn, M., Tanaka, T., “The System BoronCarbon-Niobium”, Research at the University of Vienna, 2000 to 2004, OEAD-Report, Nov. 2004, unpublished work, 1–39 (2004) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, #, 27)

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Paderno, V., Paderno, Yu., Filippov, V., Liashchenko, A., “Directional Crystallization of B4C-NbB2 and B4C-MoB2 Eutectic Compositions”, J. Solid State Chem., 117, 523–528 (2004) (Morphology, Phase Relations, Experimental, *, 14) Nunes, C.A., Kaczorowski, D., Rogl, P., Baldissera, M.R., Suzuki, P.A., Coelho, G.C., Grytsiv, A., Andre, G., Bouree, F., Okada, S., “The NbB2-Phase Revisited: Homogeneity Range, Defect Structure, Superconductivity”, Acta Mater., 53, 3679–3687 (2005) (Crys. Structure, Experimental, Electr. Prop., *, 33) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg, Germany, August 21–26, 142 (2005) (Phase Relations, Experimental, *, 4) Livramento, V., Marques, M.T., Correia, J.B., Almeida, A., Vilar, R., “Dispersion-Strengthened Nanocomposites Prepared with Niobium Carbide and Niobium Boride by Mechanical Alloying”, Mater. Sci. Forum, 514-516(Pt. 1, Advanced Materials Forum III), 707–711 (2006) (Morphology, Experimental, Mechan. Prop., 10) Tsuchida, T., Kakuta, K., “Fabrication of SPS Compacts from NbC-NbB2 Powder Mixtures Synthesized by the MA-SHS in Air Process” J. Alloys Compd., 415(1-2), 156–161 (2006) and ibid 398(1-2), 67–73 (2005) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 14) Tsuchida, T., Kakuta, K., “MA-SHS of NbC and NbB2 in Air from the Nb/B/C Powder Mixtures”, J. Eur. Ceram. Soc., 27(2-3), 527–530 (2007) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 7) Pecanha, R.M., Ferreira, F., Coelho, G.C., Nunes, C.A., Sundman, B., “Thermodynamic Modeling of the Nb-B System”, Intermetallics, 15, 999–1005 (2007) (Thermodyn., Phase Diagram, Phase Relations, 31) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)


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Boron – Carbon – Silicon Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Kostyantyn Korniyenko

Introduction Phase relations in the B-C-Si system are of great interest above all because boron, carbon and silicon are the basic elements for the development of technically important refractory ceramics and hard materials [2002Sei]. The attention of investigators of the constitution of the ternary B-C-Si system has been concentrated mainly on understanding the sintering mechanisms of SiC with boron in combination with carbon and the sintering of boron carbide with silicon [1990Tel]. Attempts of the construction of the phase diagram has been carried out by [1972Gug, 1972Kie] (partial liquidus surface projection, schematic isothermal sections as well as a series of temperature-composition sections) and by [1964Sec, 1969Sha] (the B4CSiC section). Publications concerning experimental studies of phase relations, crystal structures and thermodynamics as well as the techniques applied are listed in Table 1. However, all the available experimental data reported in literature are quite contradictory and insufficient to produce a clear interpretation of the phase relationships in the system. With a view to solving this problem, thermodynamic calculations involving the published experimental data were undertaken by [1982Doe, 1986Lil, 1994Gou1, 1994Gou2, 1995Gou, 1996Kas, 2002Sei] but the results obtained are still in need of further experimental verification. The thermodynamic properties, namely related to the vaporization behavior of B-C-Si alloys, as well as a thermodynamic analysis of the boron dissolution in silicon carbide were obtained experimentally by [1964Ver] and by [1965Mee1, 1977Saf], respectively. Reviews of the literature data relating to phase equilibria in the ternary are presented in [1972Kie, 1983Sch, 1996Kas, 2002Sei]. A review of the thermodynamic properties has been provided by [1996Sin].

Binary Systems The accepted B-C boundary binary system comes from a thermodynamic assessment and modelling carried out by [1996Kas]. A thermodynamic description for the system also appears in [1998Kas], the main differences between the two being in the modelling of the ‘B4C’ and (βB) phases, but for the purposes of this assessment, the former is accepted as the fit to the experimentally determined eutectic temperature is superior. The binary boundary B-Si system shown in Fig. 1 is accepted from [1998Fri] (based on an original publication in 1991 [1991Fri] and reproduced in [1995Lim]). The C-Si boundary binary system (Fig. 2) is accepted according to [1996Gro, 1998Gro].


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Solid Phases Crystallographic data of the solid B-C-Si phases, and their concentration and temperature ranges of stability are presented in Table 2. The joint solubility of two components in boron, carbon or silicon has not been reported. Ternary compounds of compositions SiB5C2 and Si2C2B3 were suggested by [1960Por] but their existence has not been confirmed in subsequent studies. [1996Lij] and [1996Lih] hot pressed mixtures of B4C, SiC and C at temperatures of 1800, 2000 and 2200˚C and a pressure of 25 MPa for 30 min under N2. From the X-ray diffraction studies, they found new ternary phases with crystal structures different from those of the phases in the binary systems. These ternary phases were denoted as P and M (labeled in Table 2 as τ1 and τ2, respectively) but as their appearance has not been confirmed by any other method of investigation it would seem that these two phases are metastable. [1996Kas] modeled the silicon solubility in ‘B4C’ by taking into account Si2-units occupying C-sites in the linear C-B-C chain as described by [1994Wer]. Structural characteristics and chemical bonds of Si-doped boron carbides were studied in [1999Xin] through calculations of different structural unit models by using a self-consistent-field discrete variation Xα method. The calculations show that the most energetically stable configuration for Si doped boron carbide occurs when the Si atom substitutes for B or C atoms at the end of boron carbide chain, and this allows occupation of interstitial sites. However, it is energetically difficult for Si to substitute B or C atom in the center of chain or in the icosahedral structural unit.

Invariant Equilibria Temperatures, reaction types and phase compositions relating to the invariant equilibria of the system are listed in Table 3; the corresponding reaction scheme is shown in Fig. 3. The data are based mainly on the review [2002Sei], which describes the Calphad assessment of [1996Kas]. However, some differences in the detail exist between the two publications. Also, two other thermodynamic calculations were published previously; [1982Doe, 1995Lim], the latter is also a Calphad assessment. But for the reaction scheme here, the data of [2002Sei] were preferred as the boundary binary systems used in the assessment agree in the most part with those accepted in the present report. At the same time, the compositions of the phases taking part in invariant equilibria as calculated by [1996Kas] are listed in Table 3, except for the data for the solid phases taking part in the reaction involving the liquid, SiBn, SiB6 and ‘B4C’ phases. This reaction is reported in [1996Kas] as transition, whereas [2002Sei] lists it as degenerate. All of the listed phase compositions in which the concentration of one of the components is zero, are marked as approximate. Existence of the quasibinary eutectic reaction L Ð βSiC + ‘B4C’ was reported at 2300 ± 20˚C by [1964Sec] following micrographic and X-ray studies. [1972Gug, 1972Kie], determined the eutectic temperature to be at 2240˚C using a similar array of techniques. However, [1969Sha] suggested that this reaction is in fact an invariant four-phase eutectic equilibrium L Ð βSiC + ‘B4C’ + (C)gr, which takes place at 2245 ± 5˚C. The discrepancies in these data result from an incorrect interpretation of the experimental data, in particular, the large scatter of the melting point values for the ‘eutectic-containing compositions’. Instead of the above ternary eutectic reaction, a transition reaction L + (C)gr Ð ‘B4C’ + βSiC was calculated by DOI: 10.1007/978-3-540-88053-0_21 ß Springer 2009

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[2002Sei] at 2295˚C. In total, three transition and three degenerate four-phase equilibria were proposed, but further experimental verification of their character is necessary.

Liquidus, Solidus and Solvus Surfaces Figure 4 presents the liquidus surface projection involving the stable phase equilibria, consistent with the reaction scheme (Fig. 3). The primary crystallization fields and isotherms are taken from the thermodynamic assessment of [1996Kas]. It can be seen that the boron and silicon solid solution regions as well as binary B-Si phases are pressed against the corresponding boundary side. The experimental data of [1972Gug] (the similar work is [1972Kie]) were used in the assessment of [1996Kas] and it was noted that calculated liquidus temperatures agree within the limits of experimental error. A liquidus surface projection was also reported in [1982Doe] but obsolete versions of the boundary binary systems were used in the calculation. A solidus surface projection is shown in Fig. 5 consistent with the reaction scheme and the compositions of the phases taking part in the invariant equilibria as calculated by [1996Kas] (Table 3).

Isothermal Sections A partial isothermal section for 2477˚C in the range of compositions adjacent to the βSiC phase (30 to 70 at.% Si, up to 1 at.% B) was proposed by [1996Kas] on the basis of thermodynamic calculations. Quite good agreement with the data of [1969Sha] in relation to the solubility of boron in βSiC at this temperature (about 0.1 mass% or 0.15 at.%) is observed. The isothermal section for 2327˚C was calculated by [1982Doe]. It is presented in Fig. 6 with slight changes relating to the constitution of the accepted binary boundary systems, in particular, with respect to the homogeneity range of the ‘B4C’ phase. The calculated isothermal section for 2227˚C taken from [2002Sei] is shown in Fig. 7. On comparing with the section for 2327˚C, it can be seen that the liquid phase field extends only from the B-Si side. The isothermal section for 2127˚C proposed by [1982Doe] is identical to the that for 2227˚C given by [2002Sei] with respect to constitution, but certain variations in the phase ranges exist. The isothermal section for 1927˚C was calculated by [1982Doe] across the whole range of composition. For boron contents up to 70 at.% (Fig. 8) it is similar to those for 2227 and 2127˚C, while phase fields containing (βB) and SiBn appear in the boron-rich corner. Because the B-Si binary system at the boron end was revised subsequent to the publication of [1982Doe], the corresponding part of the section is omitted in Fig. 8. [1982Doe] predicted similar phase equilibria taking place at 1727˚C. Schematic isothermal sections at 1900˚C and 1700˚C proposed by [1972Kie] are in good agreement with the calculated results of [1982Doe] for similar temperatures. [1996Lij] investigated phase equilibria of composites of approximate composition Si6.2B38.1C55.7 (in at.%) from the C-B4C-SiC region of the system. The samples were prepared by hot pressing at temperatures of 1800, 2000 and 2200˚C. According to X-ray diffraction data, the phases (C)gr, ‘B4C’ and βSiC were found in the specimens prepared at all sintering temperatures but at 2200 and 2000˚C the τ1 or τ2 phases, respectively (Table 2), were also identified. These phases seem to be metastable because they were not detected by TEM. Landolt‐Bo¨rnstein New Series IV/11E1

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Isothermal sections at 1227˚C and 1127˚C were calculated by [1996Kas] and [2002Sei], respectively. The character of phase relationships at both these temperatures is similar. In particular in the B rich corner, phase equilibria involving SiB3 appear, in contrast to higher temperatures. The section at 1127˚C according to [2002Sei] is shown in Fig. 9. [1994Gou1, 1994Gou2] presented a thermodynamic evaluation of the isothermal section at 1127˚C, taking into account their own results of phase codeposition using the chemical vapor deposition (CVD) technique. However, the character of the proposed equilibria differ from those reported in [2002Sei], and moreover show the section only schematically. As [2002Sei] presents the phase diagram data at various temperatures and these data are in a good mutual agreement, they should be considered to be more reliable.

Temperature – Composition Sections As was noted in the section “Invariant Equilibria”, the quasibinary reaction L Ð βSiC + ‘B4C’ along the SiC-B4C section was reported on the basis of experimental data by [1964Sec] (commented in [1964Sch, 1965Sec]) and by [1972Gug, 1972Kie] at 2300 ± 20˚C or at 2240˚C, respectively. At the same time, [1969Sha] observed an invariant four-phase equilibrium L Ð βSiC + ‘B4C’ + (C)gr at 2245 ± 5˚C. Thermodynamic calculation of the SiC-B4C temperature-composition section was carried out by [1982Doe], and later by [1996Kas] for the SiC-B4.6C section. This change is related with correction of the ‘B4C’ phase composition corresponding to the congruent melting point as appearing in the later version of the B-C boundary binary system. The calculated SiC-B4.6C temperature-composition section is shown in Fig. 10 according to [1996Kas, 2002Sei], accepted in this report as preferable. Properties of the SiC-B4C composites are presented in [2002Gun, 2003Aka, 2004Ueh, 2005Lee, 2005Tka, 2007Lat], the SiC/B4C multilayer coatings were studied by [1996Her]. The calculated temperature-composition sections SiC-B and the isopleth at 80 at.% B (both according to [1996Kas]) as well as the Si-B4.18C [2002Sei] are presented in Figs 11, 12 and 13, respectively. All the section are brought into conformity with accepted boundary binary systems and the reaction scheme.

Thermodynamics The vaporization behavior was studied by [1964Ver] using a mass spectrometer. The gaseous molecule SiBC was identified on the basis of measurements of mass, isotopic distribution, intensity profile in the molecular beam and appearance potential. The determined reaction enthalpies are listed in Table 4. The values of the free energy functions derived from the usual statistical thermodynamic formulas for the SiBC molecule at the temperatures of 1727, 1827, 1927 and 2027˚C as well as it’s atomization energy ΔH˚0 (at.) calculated using the third-law method are presented in Table 5. The atomization energy value demonstrates that this molecule is very strongly bonded. The partial pressures are listed in Table 6. Thermodynamic behavior of the gaseous phase during the alloying of silicon carbide with boron were investigated by [1977Saf]. The SiC single-crystals alloyed with boron were obtained by a diffusional annealing technique in the temperature range 1600 to 2550˚C. The partial heat of dissolution of boron in silicon carbide was reported as 415.9 kJ·mol–1. Thermodynamic aspects of the dissolution process of the acceptor impurity boron in silicon DOI: 10.1007/978-3-540-88053-0_21 ß Springer 2009

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carbide are also reported in [1986Lil]; the temperature dependences of the partial pressures of the interacting components were calculated for the temperature interval 1527 to 2727˚C. The calculation of the Fermi level in p-SiC (B) also was carried out. Thermodynamic calculations of phase equilibria were carried out by in [1982Doe, 1996Kas, 2002Sei]; the work of [1996Kas] is a Calphad assessment of the system, which was reported in the review by [2002Sei]. The thermodynamic description of [1996Kas] is based on data for the pure elements as stored in the SGTE database [SGTE]. Datasets for the boundary binary systems were taken from [1991Fri, 1996Gro, 1996Kas]. Because of a lack of data and the probably small energetic differences out of measurable quantities, a single analytical Gibbsenergy description was used to describe the α- and β- modifications of SiC [1996Gro]. The review by [2002Sei] reports the ternary assessment by [1996Kas], although some of the details of the calculation are slightly different to the original work. This may be owing to different unary data being used in the calculations given by [2002Sei], but no indication of this is actually given. The calculated elements of phase diagram are described in the relevant parts of the present report. The results of an experimental study and calculation of the thermodynamic parameters for chemical vapor codeposition (CVD) in a hot-wall reactor of alloys of the B-C-Si system are reported in [1994Gou1, 1994Gou2] (for 1127˚C) and in [1995Gou] (for the temperature range 927 to 1127˚C). The initial gaseous mixture consisted of methyltrichlorosilane, boron trichloride and hydrogen. The codeposits were first obtained on graphite, but the final purpose was to extend the process to composite materials. A total pressure was 0.395 bar was applied, with a total flow rate was between 0.1 and 1.0 g·min–1. It was concluded in [1995Gou] that low temperatures and high mass flow rates favor the deposition of uniform coatings but different results can be obtained as a function of the inlet gaseous composition since different kinetic limitations can arise which favor a boron or a silicon excess in the coating. A theoretical thermodynamic analysis of the SiC-B4C section [1996Sin] indicated that the adiabatic temperatures in combustion synthesis can be reduced significantly by the addition of filler (SiC or ‘B4C’ phases). The ‘B4C’ phase was shown to be much more effective in reducing the adiabatic temperature than SiC. In their opinion, these parameters can be very helpful in the selection of optimum processing conditions for the synthesis and densification of SiC-B4C composites.

Notes on Materials Properties and Applications B-C-Si alloys are of great practical interest due to their excellent mechanical properties (hardness, strength, etc.), chemical stability, high-temperature stability as well as high semiconducting properties. Among the phases taking part in equilibria in this system, silicon carbide and boron carbide are the most promising materials for use in different fields of technology. The addition of boron to silicon carbide increases its hardness, heat resistance and polishing ability at the same time as retaining the high oxidation resistance. On the other hand, the addition of silicon to boron carbide improves its sintering behavior and mechanical properties [1978Ekb]. The ceramic eutectic SiC-B4C is of interest [1979Hon] because of the improved mechanical properties and high temperature stability. These constituents also serve as the basis for engineered multilayer coatings (SiC/B4C) offering significant potential for improved tribological properties [1996Her] as well as for heterojunction diodes owing to their electronic properties [2001Ade]. Landolt‐Bo¨rnstein New Series IV/11E1

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The experimental techniques applied and properties studied are listed in Table 7. In a review of the mechanical properties of nanostructured superhard materials [2001And] a value of Vickers microhardness of 63 GPa was cited for a Si0.35B12C2.9 monolayer film [1998Bad].

Miscellaneous Investigations of the scaling resistance of B-C-Si alloys with a boron content of 80 at.% at 1000˚C carried out by [1965Mee2] show that the kinetics of oxidation of these alloys exhibits a parabolic diffusional character with a gradual decrease in the oxidation rate and the formation of a thick and strongly adhering glasslike film of the borosilicide type. According to the findings of [1977Gor], investigations of the absorption spectra associated with the excitation of excitons localized on neutral boron in the silicon carbides αSiC and α’’SiC make it possible to discover a series of bands that arise owing to the unequivalent positions of the substitutional impurities in the crystal lattice. A model of the energy spectrum of these excitations is proposed that takes into account valley-orbit interaction and explains the polarization properties of the spectra. It was reported by [1979Hon], that directionally solidified the SiC-B4C eutectic forms lamellar microstructures with no colonies if the solidification rate is below 0.02 m·h–1. [1979Pan] have studied the kinetics of wetting of ‘B4C’ by silicon. Formation of the SiC-based phase was reported. The morphology of the in-situ growth of SiC in the carbonbased composite bodies consisting of βSiC and ‘B4C’ particles was studied in [1994Oga]. A great part of SiC grains was found to be disks showing irregular utlines and rugged surface. Rod shaped grains of SiC were also observed in the fracture surface of the composites, of which cross sections revealed various morphologies such as triangular, rectangular and irregular. A semiempirical analysis of the sizes and signs of the small boron hyperfine interaction constants as determined previously by [1995Adr] using electron nuclear double resonance in boron-doped αSiC shows that the B–Si-C+ model of this centre can account for the highly unusual feature of small isotropic and anisotropic boron hyperfine constants that are positive and negative, respectively. A high-frequency pulsed EPR/ENDOR study at 95 GHz and on the 13 C-enriched single crystals of boron-doped αSiC was reported by [1997Sch]. This study enabled the formulation of a consistent model of the electronic structure of the shallow and deep boron acceptor at low temperatures. With regard to the shallow boron acceptor, it was concluded that about 40% of the spin density was located in the pz-orbital of the carbon that is nearest to boron. Later, a high-frequency (95 GHz) and conventional-frequency (9.3 GHz) pulsed EPR/ENDOR study of the deep boron acceptor in αSiC was presented in [1998Dui]. The results suggest a model in which the deep boron acceptor consists of a boron atom in a silicon position with an adjacent carbon vacancy. EPR spectra of deep boron in βSiC and αSiC crystals have been measured and studied by [1998Bar]. An ENDOR investigation has been performed on the shallow boron acceptor in the βSiC by [1998Hof]. Hyperfine and quadrupole parameters were determined with high precision for both isotopes (10B and 11B). The oxidation behavior of B4C-SiC/C composites of various compositions at temperatures up to 1500˚C was studied in [1999Guo, 2003Fan, 2003Nar], and the results indicated that the composites exhibited variable oxidation resistance at high temperatures depending on composition and oxidation temperature. The variance of self-healing properties was attributed to

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the difference in the compositions, and the properties of the decarbonized layers including wetting ability, viscosity, volatility and oxygen permeability. The sintering of nano crystalline αSiC by doping with boron carbide was carried out in [2002Dat]. The maximum density of the sintered αSiC was obtained at a concentration of 0.5 mass% of boron carbide with 1 mass % of carbon, sintered at 2050˚C for 15 min under vacuum (3 mbar). It was reported that carbon reduced the silica layer and enhanced the bulk self diffusion coefficient of silicon carbide by many orders of magnitude. It also inhibits grain growth of SiC crystals. The optimum carbon content which gives rise to high density was found to be 1 mass%. According to [2002Gun], the mass gain of directionally solidified B4C-SiC composites at 750˚C owing to oxidation is about 1/3 to 1/5 less than that of monolithic B4C. The surface perpendicular to the growth direction showed slightly better oxidation resistance than that parallel to the growth direction. The diffusion of B impurities in βSiC was analyzed by [2002Rur] using firstprinciples electronic structure calculations. Through molecular dynamics, it was found that substitutional B at a Si lattice site is readily displaced by a nearby Si interstitial by the process known as a kick-out mechanism, in agreement with recent experimental results. This is in contrast to the situation in Si, where B has recently been shown to diffuse via a mechanism involving interstitial lattice sites. The kinetics of βSiC sintering in the presence of boron at 2150˚C was studied in [2003Sto]. It was found that for a carbon concentration of 3 mass%, the optimum addition of boron was 0.2 to 0.5 mass%. The flux growth of SiC crystals from a SiCB4C eutectic melt was carried out by [2004Epe] at temperatures 2300–2350˚C. Both selfnucleated SiC platelets of up to 5 mm in diameter and epitaxial layers grown by physical vapor transport (PVT) on αSiC substrates have been produced. αSiC crystals with low boron contents have been obtained by the same technique by [2006Fan]. The effects of growth conditions, diffusion barrier coatings and hot zone materials on B incorporation into the αSiC crystals were evaluated. Following electron-microscope studies, the processes and mechanisms of structural and phase transformations in the diamond particle Si-B4C contact zone involved during sintering of a composite having an initial composition of ‘B4C’ + diamond + (αSi) are discussed in [2006Shu]. It has been shown that as the sintering temperature increases in the range from 1300 to 2000˚C, the formation of the zone microstructure is determined by the following sequence of the processes: the formation of the nano-dispersed secondary boron carbide on the diamond particle surface, thickening of this layer through the growth of boron carbide grains caused by the boron diffusion from the matrix component to the diamond particle surface, transformation of the resultant layer of anisometric grains to form at the first stage a fine-grained sub-layer in contact with diamond, which later undergoes a complete reconstruction to form a boron carbide-silicon carbide fine-grained composition. Molecular dynamics simulation of the amorphous B-C-Si system was carried out by [2006Ye] in order to investigate the diffusion behavior and analyze the influence of the addition of B on the thermal stability and creep resistance at high temperature of the amorphous system. The results show that the self-diffusion of boron tends to ascend apparently up to 1800˚C, which accounts for phase separation of the amorphous state would take place at about 1800˚C. Below this temperature, the B-C-Si system will retain the thermal stability and good creep resistance. An investigation of the chemical vapor deposition (CVD) process for the B-C-Si system is presented in [2007Ber]. By adding a porous substrate with a high internal surface in the hot zone of the reactor, the consumption of specific species is enhanced, revealing the effective precursors of the solid. In order to better understand the mechanisms of the solid formation,


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correlations are indicated between the gas phase analysis, the deposition kinetics and the deposit physicochemical characteristics. Two chemical processes (at 800–900 and at 900–1000˚C) compete against each other.

. Table 1 Investigations of the B-C-Si Phase Relations, Structures and Thermodynamics Temperature / Composition / Phase Range Studied

Reference


[1960Por]

Sintering by hot pressing at 1700–2300˚C; polishing; chemical etching; optical microscopy; X-ray diffraction; melting points measurements (Samsonov and Petrash’s method [1955Sam])

[1964Sec]

Cold and hydrostatic pressing; firing and cooling; 2250˚C, 2300˚C, 2350˚C; the SiCmelting points measurements (optical B4C section pyrometry); X-ray diffraction; optical microscopy; chemical analysis

[1964Ver]

Knudsen cell mass spectrometry

The Si-SiC-B4C-B partial system, the Si-B4C section

1507–2227˚C

[1965Mee1] Precipitation from the gaseous phase; X-ray diffraction; optical microscopy; chemical analysis

50 at.% C

[1965Mee2] Sintering by hot pressing at 1700–2000˚C; X-ray diffraction; optical microscopy; chemical analysis

80 at.% B

[1967Dok]

Hot pressing; arc melting; step annealing at The Si-SiC-B4C-B partial system 1660–1100˚C; chemical etching; X-ray diffraction; optical microscopy; chemical analysis

[1969Nie]

Reduction by hydrogen; precipitation from the gaseous phase; X-ray diffraction; optical microscopy; chemical analysis

SiB5C2, Si2C2B3

[1969Sha]

Crystal growth; firing and quenching; X-ray diffraction; optical microscopy

2200–2550˚C; the Si-B4C-C partial system

[1970Sha]

Recrystallization; crystal growth; X-ray diffraction; emission spectroscopy

αSiC with boron additions

[1971Kal]

Heat pressing; sintering; chemical etching; X-ray diffraction; optical microscopy

SiC-B section adjoining range of compositions

[1972Gug]

Hot pressing and sintering; X-ray diffraction; chemical analysis; optical microscopy; fusion; direct observation of melting in a furnace

1700–2300˚C, whole range of compositions

[1972Kie]

Hot pressing and sintering; X-ray diffraction; chemical analysis; optical microscopy; fusion; direct observation of melting in a furnace

1700–2300˚C, whole range of compositions

[1975Bin]

Hot pressing and sintering; SEM; X-ray diffraction 1750–1950˚C; (SiC) with 1 mass% of (‘B4C’) additions

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B–C–Si

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[1975Kal]

Hot pressing; sintering; X-ray diffraction; optical microscopy

SiC-B section adjoining range of compositions

[1977Saf]

Diffusional annealing; neutron-activation analysis 1600–2550˚C

[1983Wal]

Polymer pyrolysis; DTA; thermogravimetric analysis (TGA); X-ray diffraction

SiC and SiC/B4C ceramics

[1986Mor]

Annealing; SEM; TEM; Auger electron spectroscopy (AES); electron energy loss spectrometry (EELS)

1400, 1500, 1600, 1800, 2000˚C; αSiC with boron additions

[1990Tel]

X-ray diffraction; thermal analysis; EELS

B-rich corner

[1994Gie]

Self-propagating high-temperature synthesis (SHS), X-ray diffraction

SiC + (0.25–20) mass% B; (Si)+SiC; ‘B4C’ + SiC; (Si) + ‘B4C’ mixtures

[1994Gou1] Chemical vapor deposition (CVD) from an initial gaseous mixture (methyltrichlorosilane + boron trichloride + hydrogen)

1127˚C

[1994Gou2] CVD from an initial gaseous mixture (methyltrichlorosilane + boron trichloride + hydrogen)

1127˚C

[1994Wer]

Optical absorption; infrared phonon spectroscopy; Raman spectroscopy

‘B4C’ with Si additions

[1995Gou]

Chemical vapor deposition (CVD); X-ray diffraction 927–1127˚C; codeposition of B, C and Si

[1996Bar]

EPR; luminescence spectroscopy; optically αSiC with boron additions detected magnetic resonance (ODMR); deep level transient spectroscopy (DLTS)

[1996Her]

Sputter deposition; grazing incidence X-ray B4C/SiC multilayer coatings scattering (GIXS); atomic force microscopy (AFM); X-ray diffraction; SEM

[1996Lih]

Hot pressing; sintering; TEM; electron diffraction

2000˚C; Si6.2B38.1C55.7 (at.%) composite

[1996Lij]

Hot pressing; sintering; TEM; energy-dispersive spectrometry (EDS)

1800˚C, 2000˚C, 2200˚C; Si6.2B38.1C55.7 (at.%) composite

[1997Gor]

Self propagating high-temperature synthesis (SHS); SEM

SiC-B4C section (0–30 mass% B)

[1997Sch]

EPR; electron nuclear double resonance (ENDOR); αSiC with boron additions

[1998Ang]

Hot pressing; carbonization; graphitization; X-ray ≤1500˚C; C/B4C/SiC composites diffraction; TGA

[1998Bar]

EPR spectroscopy


βSiC and αSiC with boron additions

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[1998Dui]

Sublimation sandwich method; EPR; ENDOR; X-ray αSiC with boron additions diffraction

[1998Hof]

EPR; ENDOR; X-ray diffraction

βSiC with boron additions

[1998Li]

Arc melting; SEM; X-ray diffraction

B4C-SiBn (n ≥ 14)-Si composites

[1999Guo]

Grinding; sintering; thermal gravimetric/ differential thermal analysis (TG/DTA); SEM

≤1500˚C; B4C-SiC/C composites

[1999Li]

Pressing; arc melting; X-ray diffraction; SEM; TEM SiB4-B4C, SiB6-B4C, SiB14-B4C sections

[2001Ade]

Plasma-enhanced chemical-vapor deposition (PECVD); X-ray diffraction

SiC-B4C heterojunction diode;

[2001Mag]

Sintering at 1950–2200˚C; SEM; optical microscopy; EDS; X-ray diffraction

αSiC-B4C composite

[2002Gun]

Isostatic pressing; floating zone method; X-ray diffraction; SEM; TEM

SiC-B4C eutectic composites

[2003Aka]

Isostatically pressing; floating zone method; X-ray SiC-B4C eutectic composites diffraction; SEM; TEM; EMPA

[2003Nar]

Mixing, pressing, arc melting; oxidation; in situ Raman spectroscopy; X-ray diffraction; SEM

800–1500˚C; 25–60 mol% SiC - B4C

[2003Shu]

Toroid-type high pressure apparatus synthesis; 7.7 MPa, 1400–1600˚C; X-ray diffraction

Si-B4C-based cermets

[2003Sto]

Ball milling; pressing; sintering; SEM; TEM

αSiC ceramics doped by boron

[2004Ole]

Infiltration of a porous compact; EMPA

Si-B4C composites

[2004Pai]

Sintering at 2000–2100˚C; X-ray diffraction; SEM

αSiC ceramics doped by boron

[2004Rog]

Arc melting; annealing with quenching; X-ray diffraction; SEM; EMPA; wavelength-dispersive spectrometry (WDS); single crystal extraction

SiBn doped by carbon

[2004Ueh]

Hot pressing; X-ray diffraction; SEM; TEM

SiC-B4C composites

[2005Lee]

Vibration milling; hot pressing; X-ray diffraction

SiC-B4C composites

[2005Tka]

Compaction of a composite powder; sintering; SiC-B4C composites hot pressing up to 2150˚C; TEM; SEM; EMPA; X-ray diffraction

[2006Hay]

Uniaxial compaction at the pressure of 25 MPa; Reaction bonded ‘B4C’ with silicon additions sintering at 1900–2100˚C; infiltration; optical microscopy; SEM; energy-dispersive spectrometry (EDS); length-dispersive spectrometry (LDS); TEM; X-ray diffraction

[2006Shu]

Sintering at 1300–2000˚C; EMPA

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Si-B4C contact zone


B–C–Si

21




[2007Ber]

CVD; Fourier transform infrared Spectroscopy (FTIR) in-situ analysis

B-C-Si ceramics

[2007Hwa]

Ball-miling; sintering at 1400˚C; polishing; X-ray powder diffraction; SEM; EDX

1400˚C; B-Si + 5 mass% B4C

[2007Lat]

Unbalanced radio-frequency (rf) magnetron sputtering; X-ray diffraction; SEM; TEM; Raman spectroscopy; Auger- electron spectroscopy

SiC and B4C monolayers; SiC/B4C multilayer coatings

[2007Mic]

Chemical vapor infiltration (CVI); Raman spectroscopy; X-ray diffraction; TEM; polarized light optical microscopy

1200–22˚C, B-C-Si microcomposites

. Table 2 Crystallographic Data of Solid Phases Phase/ Temperature Range [˚C] (βB) < 2092

(C) gr < 3827 (sublimation point)



hR333 R 3m βB

a = 1093.30 c = 2382.52

hP4 P63/mmc C (graphite)

a = 246.12 c = 670.90

Comments/References [1993Wer] Dissolves up to 1.5 at.% C at 2098˚C [1996Kas] Dissolves up to 2.1 at.% Si at 2037˚C [1998Fri] [Mas2]

Dissolves up to 2.3 at.% B at 2382˚C [1996Kas]

(αSi) (I) cF8 < 1414 at 1.013 bar Fd 3m C (diamond)

a = 543.06

at 25˚C [Mas2, V-C2] Dissolves up to 1.1 at.% B at 1384.5˚ C [1998Fri] Dissolves up to 0.7 at.% C at 1270˚C [1998Fri]

(βSi) (II) > 9.624 bar

tI4 I41/amd βSn

a = 468.6 c = 258.5

at 25˚C [Mas2, V-C2]

(γSi) (III) > 16.208 bar

cI16 Im 3m γSi

a = 663.6

at 25˚C [Mas2, V-C2]


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(δSi) hP4 > 16.208 - 1.013 bar P63/mmc αLa

Lattice Parameters [pm] a = 380 c = 628

Comments/References at 25˚C [Mas2, V-C2]

‘B4C’ < 2458

hR45 R 3m B13C2

from 8.9 at.% B at 2098˚C to 18.9 at.% B at 2382˚C [1996Kas] a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5 labeled as B4+δC [2002Sei]; dissolves 4.3 at.% Si at 2000˚C [1972Kie] a = 560.7 in a Si10B80C10 (at.%) alloy, hotc = 1220.6 pressed [1965Mee2] a = 558 to 561 in B4C-SiC alloys hot-pressed and c = 1199 to 1211 sintered at 1800–2100˚C [1972Kie] a = 555.92; 556.37 reaction bonded ‘B4C’ [2006Hay] c = 1205.2; 1233.8 a = 563.8 in a Si0.5B13C2.5 specimen obtained c = 1231.5 by SHS [1994Gie] a = 560.03 in a Si6.2B38.1C55.7 (at.%) composite c = 1208.0 sintered at 2000˚C [1996Lih]

SiB3 < 1270

hR42 R 3m B6P

a = 631.9 c = 1271.3

[V-C2] 73 to 74 at.% B [1998Fri]

SiB6 < 1850

oP340 Pnnm SiB6

a = 1439.7 b = 1831.8 c = 991.1

[V-C2] 85.4 to 86.2 at.% B [1998Fri]

SiBn < 2037

hR36 R 3m βB hR339 R 3m FeB49

SiBnCx βSiC < 2824

[Mas2], n ≈ 23; 94.1 to 98.5 at.% B [1998Fri] a = 1101 c = 2390

[V-C2]

a = 1101.52 c = 2386.25

n = 30.4, x = 0.35, single crystal data [2004Rog]

cF8 a = 435.81 F 43m ZnS (sphalerite)

a = 435.3

a = 436

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50 at.% C [V-C2] labeled elsewhere as SiC of the 3C type dissolves 16 at.% B at 1800˚C [1972Kie] dissolves 3 at.% B (SHS obtained specimens) [1994Gie] in a Si30B20C50 (at.%) alloy, T = 1600˚C [1965Mee1]; Si50C50 (at.%) specimen sintered at T = 1950˚C in the SiC-B4C composite sintered at 1500˚C [2002Gun] Landolt‐Bo¨rnstein New Series IV/11E1

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B–C–Si

. Table 2 (continued) Phase/ Temperature Range [˚C] αSiC

Pearson Symbol/ Space Group/ Prototype hP12 P63mc αSiC

Lattice Parameters [pm] a = 308.07 c = 1511.74

Comments/References 50 at.% C [V-C2] labeled elsewhere as SiC of the 6H type, also known as moissanite; metastable; dissolves 0.2 mass% B at 2450–2500˚C [1970Sha] in the Si6.2B38.1C55.7 (at.%) composite sintered at 2000˚C [1996Lih]

a = 307.3 c = 1508 α’SiC

hP42 P 3m1 α’SiC

a = 307 c = 528.7

50 at.% C [V-C2] metastable; tentative structure

α’’SiC

hR48 R 3m α’’SiC

a = 308.2 c = 604.9

50 at.% C [V-C2] labeled elsewhere as SiC of the 15R type; metastable

α’’’SiC

hR186 R 3m α’’’SiC

a = 307 c = 2341.7

50 at.% C [V-C2] labeled elsewhere as SiC of the 15R type; tentative structure; metastable

τ1

o**

a = 850 b = 900 c = 490

in the Si6.2B38.1C55.7 (at.%) composite sintered at 2200˚C, together with the (C)gr, ‘B4C’ and βSiC phases [1996Lij]; metastable?

τ2

m**

a = 900 b = 590 c = 540 β = 119.3˚

in the Si6.2B38.1C55.7 (at.%) composite sintered at 2000˚C, together with the (C)gr, ‘B4C’ and βSiC phases; labelled as M [1996Lih]; metastable?

. Table 3 Invariant Equilibria Composition (at.%) Reaction L + (C)gr Ð ‘B4C’ + βSiC


T [˚C] 2295

Type

Phase

U1

L

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B

C

Si

61.7

25.0

13.3

(C)gr

2.0

98.0

0.0

‘B4C’

81.0

19.0

0.0

βSiC

0.1

49.9

50.0

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. Table 3 (continued) Composition (at.%) T [˚C]

Reaction L + (βB) Ð ‘B4C’ + SiBn

2005

Type

Phase

B

C

Si

U2

L

93.0

0.2

6.8

(βB)

98.4

1.4

0.2

‘B4C’

91.0

8.9

0.1

SiBn

96.8

0.0

3.2

L + SiBn Ð SiB6, ‘B4C’

1850

D1

L

62.1

0.0

37.8

L + βSiC Ð ‘B4C’ + (αSi)

1396

U3

L

L Ð (αSi) + SiB6, ‘B4C’

(αSi) + SiB6 Ð SiB3, ‘B4C’

1384

1198

D2

D3

4.7

0.0

95.3

βSiC

0.7

49.3

50.0

‘B4C’

85.8

14.0

0.2

(αSi)

0.6

0.0

99.4

L

8.1

0.0

91.9

(αSi)

0.9

0.0

99.1

SiB6

85.4

0.0

14.6

‘B4C’

86.3

12.6

1.1

(αSi)

85.5

0.0

14.5

SiB6

0.5

0.0

99.5

SiB3

73.8

0.0

26.2

‘B4C’

86.8

13.1

0.1

. Table 4 Thermodynamic Data of Reaction or Transformation Quantity, per mole of atoms [kJ, mol, K]

Reaction or Transformation

Temperature [˚C]

1/3{SiBC + SiC2 Ð BC2 + Si2C}

1507–2227˚C 36.8 ± 5.9

1/3{SiBC + Si Ð Si2C + B}

1507–2227˚C –37.6 ± 7.1

1/3{SiBC + Si Ð Si2 + B + (C)}

1507–2227˚C 12.5 ± 6.7

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Comments [1964Ver] Knudsen effusion. The state of the species is gas.


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. Table 5 Thermodynamic Properties of Single Phases Phase

Temperature Range [˚C]

Property, per mole of atoms [J, mol, K]

SiBC

1727

–{(GT0 – H˚0)/T} = 298.9

1827 1927 2027

–{(G

T

– H˚0)/T} = 301.4

–{(G

T

– H˚0)/T} = 303.9

–{(G

T

– H˚0)/T} = 306.4

0 0 0

Comments [1964Ver] Knudsen effusion. The state of the species is gas.

ΔH˚0 (at.) = 1032460 ± 25100

. Table 6 Vapor Pressure Measurements Phase(s) SiBC

Temperature [˚C]

Pressure [bar]

Comments

1976

log10 (pSiBC) = –6.62

[1964Ver] Knudsen effusion.

2032

log10 (pSiBC) = –6.34

The state of the species is gas.

2071

log10 (pSiBC) = –6.11

2028

log10 (pSiBC) = –6.77

1877

log10 (pSiBC) = –6.77

1893

log10 (pSiBC) = –6.46

1810

log10 (pSiBC) = –6.66

. Table 7 Investigations of the B-C-Si Materials Properties Reference [1960Por]


Type of Property

Hydrostatic measurements; pycnometry; Specific gravity; thermo-emf; mechanical and electrical properties tests microhardness; electrical resistivity

[1965Mee1] Microhardness measurements

Microhardness

[1965Mee2] Microhardness and electrical resistivity measurements

Microhardness; specific electrical resistivity

[1967Dok]

PMT-3 microhardness tests

Microhardness

[1969Nie]

Mechanical properties tests

Microhardness

[1971Kal] [1975Kal]

PMT-3 microhardness tests; electrical resistivity and density measurements

Microhardness; specific electrical resistivity; density

[1972Gug] [1972Kie]

Vickers microhardness tests

Microhardness


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. Table 7 (continued) Reference [1975Bin]


Type of Property

Sintering under pressure

Densification behavior

[1978Ekb]

Strength, Vickers microhardness tests

Strength, microhardness

[1979Hon]

Microhardness tests (Leitz microhardness Microhardness tester, Knoop diamond indenter)

[1979Pan]


Microhardness

[1983Wal]

Strength, Vickers microhardness tests

Flexural strength, microhardness

[1987Bou]

Microhardness tests (Leitz microhardness Microhardness; Young’s modulus; tester, Knoop diamond indenter); ductility dynamic method (Grindo Sonic); point method

[1994Wer]

Optical absorption; infrared phonon spectroscopy; Raman spectroscopy

Absorption coefficient; reflectivity; absorption index; Raman intensity

[1995Adr]

Electron nuclear double resonance

Boron acceptor behavior

[1995Mas]

PMT-3 mechanical tests

Microhardness; microbrittleness; microstrength; crack resistance; density

[1996Bar]

EPR spectroscopy (X-band and Q-band spectrometer); photoluminescence (PL) spectroscopy

EPR and PL spectra

[1996Her]

Double crystal diffraction topography (DCDT); grazing incidence X-ray scattering (GIXS); atomic force microscopy (AFM); SEM

Residual stress; structure of films

[1997Gor]

Electrical resistivity and strength measurements

Electrical resistivity; bending strength

[1997Shi]

Strength, electrical resistance, thermal emf measurements

Compressive strength, specific electrical resistance, the coefficient of thermal emf, volt-ampere characteristic; density

[1998Ang]

Thermogravimetric analysis (TGA) (Seiko SSC 5200 apparatus); thermomechanical analysis (TMA) (Perkin-Emler thermal analyzer); density and porosity tests; Knoop hardness measurements; 4 point bend test

Density; hardness; flexural strength; flexural modulus; porosity

[1998Bad]


Microhardness

[1998Bar]

EPR spectroscopy

EPR spectra

[1998Dui]

EPR spectroscopy; electron-nuclear EPR, ENDOR, ESEEM spectra double resonance (ENDOR); electron spin echo envelope modulation (ESEEM)

[1998Li]

Four-probe method; ac method; laser-flash technique

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Electrical conductivity; the Seebeck coefficient; thermal conductivity Landolt‐Bo¨rnstein New Series IV/11E1

B–C–Si

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Type of Property

[1999Li]

Four-probe method; ac method; laserflash technique; Archimede’s method

[2000Gri]

Indentation and SENVB methods; three- Strength; Young’s modulus; hardness; point bending test; measuring the rate of crack resistance; thermal stress propagation of an acoustic signal; Vickers hardness tests

[2001Ade]

Current and voltage measurements (Keithley 2400-C SourceMeter and TestPoint software)

[2001Mag]

Archimede’s method; four-point bending Density; flexural strength; Weibull test; hardness tests (Leitz hardness tester) modulus; hardness; fracture toughness

[2002Dar]

Tensile creep tests (PSB 100 Hydropuls Schenck machine); SEM; TEM

Creep curves

[2002Gun]

DC four-probe method; laser-flash technique; Vickers microhardness tests

Electrical conductivity; thermal conductivity; microhardness

[2002Shi]

Microhardness and electrical conductivity Microhardness; electrical conductivity measurements

[2003Shu]

Vickers indentation method (PMT-3 tester)

Microhardness; hardness; fracture toughness

[2004Pai]

DC four-probe method; laser-flash technique

Electrical conductivity; thermoelectric power; Seebeck coefficient; thermal conductivity; thermal diffusivity; density

[2004Ueh]

Vickers indentation method; voltage measurements

Hardness; fracture toughness; Seebeck coefficient

[2005Lee]

Vickers indentation method; wear tests (reciprocating ball-on-disk tester TE77)

Hardness; wear resistance

[2005Tka]

Mechanical properties tests

Microhardness; cracking resistance; bending strength; density

[2006Shu]

Mechanical tests

Strength

[2006Ye]

Molecular dynamics simulation; mechanical tests

Thermal stability; creep resistance

[2007Hwa]

Vickers hardness measurements

Hardness

[2007Lat]

Vickers indentation method; scratch test Hardness; residual stress; critical load of (CSEM-Revetest) Failure; Young’s modulus

[2007Mic]

Tensile tests; strain measurements (compliance calibration technique);


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Electrical conductivity; the Seebeck coefficient; thermal conductivity; density

Current-voltage curves

Young’s modulus; thermal expansion; microcomposite’s failure stress; fracture behavior

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. Fig. 1 B-C-Si. The B-Si phase diagram

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. Fig. 2 B-C-Si. The C-Si phase diagram


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B–C–Si

. Fig. 3 B-C-Si. Reaction scheme

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. Fig. 4 B-C-Si. Calculated liquidus surface projection


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B–C–Si

. Fig. 5 B-C-Si. Calculated solidus surface projection

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. Fig. 6 B-C-Si. Calculated isothermal section at 2327˚C


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. Fig. 8 B-C-Si. Calculated partial isothermal section at 1927˚C


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. Fig. 10 B-C-Si. Calculated temperature - composition section SiC-B4.6C


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. Fig. 11 B-C-Si. Temperature - composition section SiC-B

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. Fig. 12 B-C-Si. Calculated temperature - composition section at 80 at.% B


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. Fig. 13 B-C-Si. Calculated temperature - composition section Si-B4.18C

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References [1955Sam]

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[1964Sec]

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[1975Bin] [1975Kal]

Samsonov, G.V., Petrash, Ye.V., “Some Physical-Chemical Properties of Titanium Boride and Nitride Alloys” (in Russian), Metalloved. Term. Obrab. Met., (4), 19–24 (1955) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Phys. Prop., 10) Portnoi, K.I., Samsonov, G.V., Solonnikova, L.A., “Alloys of the Boron-Silicon-Carbon System”, Russ. J. Inorg. Chem. (Engl. Transl.), 5(9), 988–993 (1960), translated from Zh. Neorg. Khim. SSSR, 5 (9), 2032–2041 (1960) (Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., Phys. Prop., 25) Schaffer, P.T.B., Hannam, A.L., “Comments on Phase Equilibria in the System Boron Carbide-Silicon Carbide by D.R. Secrist”, J. Am. Ceram. Soc., 47(11), 594–595 (1964) (Crys. Structure, Phase Relations, Assessment, 9) Secrist, D.R., “Phase Equilibria in the System Boron Carbide-Silicon Carbide”, J. Am. Ceram. Soc., 47(3), 127–130 (1964) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 3) Verhaegen, G., Stafford, F.E., Drowart, J., “Mass Spectrometric Study of the Systems Boron-Carbon and Boron-Carbon-Silicon”, J. Chem. Phys., 40(6), 1622–1628 (1964) (Thermodyn., Experimental, 31) Meerson, G.A., Kiparisov, S.S., Gurevich, M.A., Fen-Sian, D., “Obtaining and Investigation of the Properties of Solid Solutions on a Pseudobinary Section SiC-BC by the Method of Joint from a Gaseous Phase” (in Russian), Poroshk. Metall., (2), 15–21 (1965) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 6) Meerson, G.A., Kiparisov, S.S., Gurevich, M.A., Fen-Sian, D., “Investigation of Conditions for Obtaining of Solid Alloys of the Pseudobinary System B4C-B4Si” (in Russian), Poroshk. Metall., (3), 62–68 (1965) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Kinetics, Mechan. Prop., 9) Secrist, D.R., “Reply to Comments on “Phase Equilibria in the System Boron Carbide-Silicon Carbide”, J. Am. Ceram. Soc., 48(4), 215–215 (1965) (Crys. Structure, Phase Relations, Assessment, 5) Dokukina, I.V., Kalinina, A.A., Sokhor, M.I., Shamrai, F.I., “On the Problem of Chemical Compounds in the System Silicon-Boron-Carbon” (in Russian), Izv. Akad. Nauk SSSR, Neorgan. Mater., 3(4), 630–637 (1967) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 8) Niemyski, T., Appenheimer, S., Panczyk, J., Badzian, A., “Vapor Phase Crystallization of B-Si-C Phase”, J. Cryst. Growth, 5(5), 401–404 (1969) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., 6) Shaffer, P.T.B., “The SiC Phase in the System SiC-B4C-C”, Mater. Res. Bull., 4(3), 213–220 (1969) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 12) Shaffer, P.T.B., “Solubility of Boron in Alpha Silicon Carbide”, Mater. Res. Bull., 5, 519–521 (1970) (Crys. Structure, Morphology, Phase Relations, Experimental, 4) Kalinina, A.A., Sokhor, M.I., Shamrai, F.I., “Investigation of the Alloys of the System Si-B-C” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 7(5), 778–785 (1971) (Crys. Structure, Morphology, Phase Relations, Experimental, Phys. Prop., 15) Gugel, E., Kieffer, R., Leimer, G., Ettmayer, P., “Investigation in the Ternary System Boron-CarbonSilicon”, Nat. Bur. Stand. Spec. Pub., Sol. State Chem., Proc. 5th Mater. Res. Symp., 364, 505–513 (1972) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Mechan. Prop., *) as quoted by [1996Kas] Kieffer, R., Gugel, E., Leimer, G., Ettmayer, P., “Investigations in the System Boron-Carbon-Silicon“ (in German), Berich. Deut. Keram. Gesel., 49(2), 41–46 (1972) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, Review, Mechan. Prop., *, 39) Bind, J.M., Biggers, J.V., “Hot-Pressing of Silicon Carbide with 1% Boron Carbide Addition”, J. Am. Ceram. Soc., 58(7-8), 304–306 (1975) (Crys. Structure, Morphology, Experimental, Phys. Prop., 19) Kalinina, A.A., Sokhor, M.I., “Phase Composition and Some Properties of Silicon-Boron-Carbon System Alloys Adjacent to the Section Slicon Carbide-Boron” (in Russian), Vysokotemperatur. Karbidy, Naukova Dumka, Kiev, 96–99 (1975) (Crys. Structure, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., Phys. Prop., 6)


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21 [1977Gor]

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[1986Mor]

[1987Bou]

[1990Ase]

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B–C–Si Gorban’, I.S., Krokhmal, A.P., “Structure of the Spectrum of Excitons Bound to Neutral Boron Atoms in Silicon Carbide”, Sov. Phys. - Solid State (Engl. Transl.), 19(5), 733–755 (1977) (Morphology, Experimental, Theory, Electronic Structure, 8) Safaraliev, G.K., Tairov, Yu.M., Tsvetkov, V.F., “Thermodynamic Analysis of Solubility and Transformation Coefficient of Boron in Silicon Carbide” in “Svoistva Legirovannykh Poluprovodnikov” (in Russian), Nauka, Moscow, 53–58 (1977) (Thermodyn., Calculation, Experimental, 6) Ekbom, L.B., “Effect of Increased Boron Content on the Sintering Behaviour and Mechanical Properties of Boron Carbide”, unknoun source, a copy is available, 183–189 (1978) (Morphology, Experimental, Mechan. Prop., 6) Hong, J.-D., Spear, K.E., Stubican, V.S., “Directional Solidification of SiC-B4C Eutectic: Growth and Some Properties”, Mater. Res. Bull., 14(6), 775–783 (1979) (Morphology, Experimental, Kinetics, Mechan. Prop., 13) Panasyuk, A.D., Oreshkin, V.D., Maslennikova, V.R., “Kinetics of the Reactions of Boron Carbide with Liquid Aluminium, Silicon, Nickel and Iron”, Sov. Powder Metall. Met. Ceram., 199(7), 487–490 (1979), translated from Poroshk. Metall., 199(7), 79–83 (1979) (Morphology, Experimental, Kinetics, Mechan. Prop., 9) Doerner, P., “Constitutional Investigations on High Temperature Ceramics of the B-Al-C-Si-N-O System by Means of Thermochemical Calculations” (in German), Thesis, University Stuttgart, Institut fuer Metallkunde, 1–194 (1982) (Phase Diagram, Phase Relations, Thermodyn., Calculation, *, 126) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Phase Diagram, Phase Relations, Review, 154) Walker, B.E., Rice, R.W., Becher, P.F., Bender, B.A., Coblenz, W.S., “Preparation and Properties of Monolithic and Composite Ceramics Produced by Polymer Pyrolysis”, Amer. Ceram. Soc. Bul., 62(8), 916–923 (1983) (Crys. Structure, Morphology, Experimental, Phys. Prop., 9) Lilov, S.K., “Thermodynamic Study of the Solubility Process of Boron in Silicon Carbide, Grown from the Vapour Phase”, J. Phys. Chem. Solids, 47(3), 245–250 (1986) (Thermodyn., Calculation, Review, Electronic Structure, 16) More, K.L., Carter, C.H., Bentley J. Jr., Wadlin, W.H., LaVanier L., Davis, R.F., “Occurrence and Distribution of Boron-Containing Phases in Sintered α-Silicon Carbide”, J. Am. Ceram. Soc., 69(9), 695–698 (1986) (Crys. Structure, Morphology, Phase Relations, Experimental, 17) Bougoin, M., The´venot, F., Dubois, J., Fantozzi, G., “Synthesis and Thermomechanical Properties of the Ceramic Dense Composites Boron Carbide-Silicium Carbide” (in French), J. Less-Common Met., 132(2), 209–228 (1987) (Morphology, Experimental, Phys. Prop., 55) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides”, Freer R. (Ed.), “The Physics and Chemistry of Carbides, Nitrides and Borides”, Kluwer Academic Publishers, Dordrecht 97–111 (1990) (Crys. Structure, Experimental, Review, 14) Telle, R., “Structure and Properties of Si-Doped Boron Carbide” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Kluwer Academic Publishers, Dordrecht (Germany), 249 (1990) (Crys. Structure, Phase Diagram, Phase Relations, Experimental) as quoted by [2002Sei] Fries, S., Lim, S-K., Lukas, H.L., “Optimisation of the Al-C, Al-Sn, B-Si, Mg-Zn, Ci-C and Zn-Sn Binary Systems as well as of the Al-Sn-Zn Ternary System” in “Leuven Proceedings, COST 507, New Light Alloys, Part C Thermodynamic Evaluation and Calculation”, Effenberg, G. (Ed.)/, /Commission of the European Communities, Part C, D9–1–5 (1991) (Phase Diagram, Thermodyn., Assessment, Calculation, 0) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of Carbondoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, Electronic Structure, 51) Gierlotka, S., Oleksyn, O., Palosz, B., “Si-C-B System Synthesized by SHS Method: Phase Diagram and Crystal Structure Evolution”, Mater. Sci. Forum, 166-169, 529–538 (1994) (Crys. Structure, Phase Relations, Experimental, 11) Goujard, S., Vandenbulcke, L., Bernard, C., “Thermodynamic Study of the Chemical Vapour Deposition in the Si-B-C-H-Cl System”, Calphad, 18(4), 369–385 (1994) (Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, *, 35)

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B–C–Si [1994Gou2]

[1994Oga] [1994Wer] [1995Adr]

[1995Gou]

[1995Lim]

[1995Mas]

[1996Bar]

[1996Gro] [1996Her]

[1996Kas]

[1996Lih]

[1996Lij]

[1996Sin] [1997Gor]

[1997Sch]

[1997Shi]

[1998Ang]

21

Goujard, S., Vandenbulcke, L., Bernard, C., Blondiaux, G., Debrun, J.L., “Thermodynamic and Experimental Study of the Chemical Vapor Codeposition in the Silicon-Boron-Carbon System at 1400 K”, J. Electrochem. Soc., 141(2), 452–460 (1994) (Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, *, 48) Ogawa, I., “Morphologies of SiC Growing in C/SiC/B4C Composites” (in Japanese), J. Ceram. Soc. Jpn., 102(8), 802–803 (1994) (Morphology, Experimental, 4) Werheit, H., Kuhlmann, U., Laux, M., “Solid Solutions of Silicon in Boron-Carbide-Type Crystals”, J. Alloys Compd., 209, 181–187 (1994) (Crys. Structure, Phase Relations, Experimental, Phys. Prop., 14) Adrian, F.J., Greulich-Weber, S., Spaeth, J.-M., “Further Evidence for the B–Si-C+ Model of the Boron Acceptor in 6H Silicon Carbide from a Theoretical Analysis of the Hyperfine Interactions”, Solid State Commun., 94(1), 41–44 (1995) (Morphology, Experimental, Electronic Structure, Phys. Prop., 9) Goujard, S., Vandenbulcke, L., Bernard, C., “On the Chemical Vapor Deposition of Si/B/C-Based Coatings in Various Conditions of Supersaturation”, J. Eur. Ceram. Soc., 15(6), 551–561 (1995) (Thermodyn., Calculation, Experimental, 19) Lim, S.K, Lukas, H.L., “Thermodynamic Optimisation of the System B-C-Si and its Boundary Systems” in “Hochleistungskeramik, Herstellung, Aufbau und Eigenschaften”, Deutsche Forschungsgemeinschaft, Petzow, G., Tobolski, J., Telle, R. (Eds.), VCH, Weinheim, 605–616, (Phase Diagram, Thermodyn., Assessment, Calculation, 57) Maslennikova, V.R., Belkina, A.A., Panasyuk, A.D., Struk, L.I., Smirnov, V.P., “Effect of High Pressure on the Structure and Properties of Materials Based on Boron Carbide”, Powder Metall. Met. Ceram., 34(9,10), 491–495 (1995), translated from Poroshk. Metall., (9/10), 3–7 (1995) (Morphology, Experimental, Mechan. Prop., Phys. Prop., 6) Baranov, P.G., Mokhov, E.N., “Electron Paramagnetic Resonance of Deep Boron in Silicon Carbide”, Semicond. Sci. Technol., 11(4), 489–494 (1996) (Crys. Structure, Experimental, Theory, Electronic Structure, Magn. Prop., Optical Prop., 25) Groebner, J., Lukas, H.L., Aldinger, F., “Thermodynamic Calculation of the Ternary System Al-Si-C”, Calphad, 20(2), 247–254 (1996) (Phase Diagram, Thermodyn., Calculation, 37) Hershberger, J., Ying, T., Kustas, F., Fehrenbacher, L., Yalisove, S.M., Bilello, J.C., “Residual Stress, Atomic Structure, and Growth Morphology in B4C/SiC Multilayer Coatings”, Surf. Coat. Technol., 86-87(1–3), 237–242 (1996) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 23) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institut, Dissertation, Stuttgart, 1–225 (1996) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calculation, Experimental, Review, #, 170) Lihua, Z., Lijun, W., Qizhong, H., Qiaoqin, Y., Shaoli, L., “Structure of C-B4C-SiC Composites with Silicon Additive”, J. Mater. Sci. Lett., 15(4), 353 (1996) (Crys. Structure, Morphology, Phase Relations, Experimental, *, 10) Lijun, W., Qizhong, H., Qiaoqin, Y., Lihu, Z., Zhongyu, X., “Effect of Sintering Temperature on Structure of C-B4C-SiC Composites with Silicon Additive”, Scr. Mater., 35(1), 123–127 (1996) (Crys. Structure, Morphology, Experimental, *, 12) Singh, M., “Thermodynamic Analysis for the Combustion Synthesis of SiC-B4C Composites”, Scr. Mater., 34(6), 923–927 (1996) (Thermodyn., Review, Theory, 12) Go´rny, G., Raczka, M., Stobierski, L., Wojnar, L., Pampuch, R., “Microstructure-Property Relationship in B4C-βSiC Materials”, Solid State Ionics, 101-103 (Pt.2), 953–958 (1997) (Crys. Structure, Morphology, Experimental, Electr. Prop., Mechan. Prop., 4) Schmidt, J., Matsumoto, T., Poluektov, O.G., Arnold, A., Ikoma, T., Baranov, P.G., Mokhov, E.N., “High-Frequency EPR Studies of Shallow and Deep Boron Acceptors in 6H-SiC”, Mater. Sci. Forum, Defects in Semiconductors-19, 258-263, 703–708 (1997) (Crys. Structure, Experimental, Theory, Electronic Structure, 33) Shipilova, L.A., Petrovskii, V.Ya., Chugunova, S.I., “Structure Formation and Electrophysical Properties of a Silicon Carbide-Boron Carbide Sintered Composite”, Powder Metall. Met. Ceram., 36(11-12), 652–656 (1997), translated from Poroshk. Metall., (11–12), 97–101 (1997) (Morphology, Experimental, Electr. Prop., Mechan. Prop., 6) Angelovici, M.M., Bryant, R.G., Northam, G.B., Roberts, A.S., Jr., “Carbon/Ceramic Microcomposites, Preparation and Properties”, Mater. Lett., 36(5-6), 254–265 (1998) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 12)


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21 [1998Bad]

[1998Bar]

[1998Dui]

[1998Fri]

[1998Gro]

[1998Hof]

[1998Kas]

[1998Li]

[1999Guo]

[1999Li]

[1999Xin]

[2000Gri]

[2001Ade]

[2001And]

[2001Mag]

[2002Dar] [2002Dat]

[2002Gun]

B–C–Si Badzian, A., Badzian, T., Drawl, W.D., Roy, R., “Synthesis and Properties of the B-C-Si and Si-N-C Hard Materials”, Diamond Related Compd., (7), 1519–1524 (1998) (Morphology, Experimental, Mechan. Prop.) as quoted by [2001And] Baranov, P.G., Il‘in, I.V., Mokhov, E.N., “Electron Paramagnetic Resonance of Deep Boron Acceptors in 4H-SiC and 3C-SiC Crystals”, Phys. Solid State, 40(1), 31–34 (1998) (Crys. Structure, Experimental, Electronic Structure, Magn. Prop., 12) von Duijn-Arnold, A., Ikoma, T., Poluektov, O.G., Baranov, P.G., Mokhov, E.N., Schmidt, J., “Electronic Structure of the Deep Boron Acceptor in Boron-Doped 6H-SiC”, Phys. Rev. B, 57(3), 1607–1619 (1998) (Crys. Structure, Experimental, Electronic Structure, Magn. Prop., 42) Fries, S., Lukas, H.L., “System B-Si” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 126–128 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, 1) Groebner, J., Lukas, H.L., Aldinger, F., “System C-Si” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 132–133 (1998) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, 1) Hofstaetter, A., Meyer, B.K., Scharmann, A., Baranov, P.G., Ilyin, I.V., Mokhov, E.N., “X-Band ENDOR of Boron and Beryllium Acceptors in Silicon Carbide”, Mater. Sci. Forum, Silicon Carbide, III-Nitrides and Related Materials, 264-268, 595–598 (1998) (Crys. Structure, Experimental, Electronic Structure, 6) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Li, J., Goto, T., Hirai, T., “Microstructure and Thermoelectric Properties of B4C-SiBn-Si Composites Prepared by Arc Melting”, J. Ceram. Soc. Jpn., 106(2), 194–197 (1998) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Phys. Prop., 16) Guo, Q., Song, J., Liu, L., Zhang, B., “Relationship Between Oxidation Resistance and Structure of B4C-SiC/C Composites with Self-Healing Properties”, Carbon, 37(1), 33–40 (1999) (Morphology, Phase Relations, Experimental, Kinetics, 22) Li, J., Goto, T., Hirai, T., “Thermoelectric Properties of B4C-SiBn (n = 4, 6, 14) In-situ Composites”, Mater. Trans., JIM, 40(4), 314–319 (1999) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Phys. Prop., 25) Xinmin, M., Cewen, N., Kefeng, C., “Structural Characteristics and Quantum Chemistry Calculation of Si-Doped Boron Carbides” in “Multiscale Modelling of Materials”, Proc. Symp. Mater. Res. Soc., Warrendale, PA, USA, 579–584 (1999) (Crys. Structure, Calculation, Electronic Structure, 12) Grigor’ev, O.N., Gogotsi, G.A., Gogotsi, Yu.G., Subbotin, V.I., Brodnikovskii, N.P., “Synthesis and Properties of Ceramics in the SiC-B4C-MeB2 System”, Powder Metall. Met. Ceram., 39(5-6), 239–250 (2000) (Morphology, Experimental, Mechan. Prop., 15) Adenwalla, S., Welsch, P., Harken, A., Brand, J.I., Sezer, A., Robertson, B.W., “Boron Carbide/n-Silicon Carbide Heterojunction Diodes”, Appl. Phys. Lett., 79(26), 4357–4359 (2001) (Crys. Structure, Experimental, Electr. Prop., 7) Andrievski, R.A., “Superhard Materials Based on Nanostructured High-Melting Point Compounds: Achievements and Perspectives”, Iner. J. Ref. Met. Hard Mater., 19(4-6), 447–452 (2001) (Morphology, Review, Mechan. Prop., 59) Magnani, G., Beltrami, G., Minoccari, G.L., Pilotti, L., “Pressureless Sintering and Properties of αSiCB4C Composite”, J. Eur. Ceram. Soc., 21(5), 633–638 (2001) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., Phys. Prop., 19) Darzens, S., Farizy, G., Vicens, J., Chermant, J.-L., “Microstructure and Creep of SiCf-SiBC”, Key Eng. Mater., 206-213(2), 989–992 (2002) (Morphology, Experimental, Mechan. Prop., 18) Datta, M.S., Bandyopadhyay, A.K., Chaudhuri, B., “Sintering of Nano Crystalline α Silicon Carbide by Doping with Boron Carbide”, Bull. Mater. Sci. (India), 25(3), 181–189 (2002) (Morphology, Experimental, Kinetics, Interface Phenomena, 22) Gunjishima, I., Akashi, T., Goto, T., “Characterization of Directionally Solidified B4C-SiC Composites Prepared by a Floating Zone Method”, Mater. Trans., 43(9), 2309–2315 (2002) (Crys. Structure, Morphology, Experimental, Electr. Prop., Kinetics, Mechan. Prop., Phys. Prop., 24)

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[2002Sei] [2002Shi]

[2003Aka]

[2003Fan]

[2003Nar]

[2003Shu]

[2003Sto] [2004Epe]

[2004Ole]

[2004Pai] [2004Rog]

[2004Ueh]

[2005Lee]

[2005Tka]

[2006Fan]

[2006Hay]

[2006Shu]

21

Rurali, R., Godignon, P., Rebollo, J., Ordejon, P., Hernandez, E., “Theoretical Evidence for the KickOut Mechanism for B Diffusion in SiC”, Appl. Phys. Lett., 81(16), 2989–2991 (2002) (Morphology, Review, Theory, Electronic Structure, Interface Phenomena, Kinetics, 20) Seifert, H.J., Aldinger, F., “Phase Equilibria in the Si-B-C-N System”, Struct. Bonding, 101, 1–58 (2002) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Calculation, Review, #, 275) Shipilova, L.A., Petrovskii, V.Ya., “Structure Formation, Electrophysical and Mechanical Properties of an Electrically Conducting Ceramic Composite Based on Silicon and Boron Carbides”, Powder Metall. Met. Ceram., 41(3-4), 147–149 (2002), translated from Poroshk. Metall., 3-4(424), 41–44 (2002) (Morphology, Experimental, Electr. Prop., Mechan. Prop., 3) Akashi, T., Gunjishima, I., Goto, T., “Characterization of Directionally Solidified B4C-TiB2 and B4CSiC Eutectic Composites Prepared by Floating-Zone Metod”, Key Eng. Mater., 247, 209–212 (2003) (Crys. Structure, Morphology, Experimental, 3) Fan, Z., Wei, T., Shi, J., Zai, G., Song, J., Liu, L., Li, J., Chen, J., “New Route for Preparation of SiCB4C/C Composite with Excellent Oxidation Resistance up to 1400˚C”, J. Mater. Sci. Lett., 22(3), 213–215 (2003) (Morphology, Experimental, Kinetics, 10) Narushima, T., Goto, T., Maruyama, M., Arashi, H., Iguchi, Y., “Oxidation of Boron Carbide-Silicon Carbide Composite at 1073 to 1773 K”, Mater. Trans., JIM, 44(3), 401–406 (2003) (Morphology, Phase Relations, Experimental, Kinetics, 33) Shulzhenko, A.A., Stratiychuk, D.A., Gargin, V.G., Belyavina, N.N., “Preparation and PhysicoMechanical Properties of B-C-Si-Based Cermet”, J. Superhard Mater., 25(5), 74–76 (2003), translated from Sverkhtverd. Mater. (Ukraine), (5), 82–84 (2003) (Morphology, Phase Relations, Experimental, Mechan. Prop., 9) cited from abstract Stobierski, L., Gubernat, A., “Sintering of Silicon Carbide. II. Effect of Boron”, Ceram. Intern., 29(4), 355–361 (2003) (Phase Relations, Morphology, Experimental, Kinetics, Transport Phenomena, 13) Epelbaum, B.M., Gurzhiyants, P.A., Herro, Z., Bickermann, M., Winnacker, A., “Flux Growth of SiC Crystals from Eutectic Melt SiC-B4C”, Mater. Sci. Forum, 457-460, 119–122 (2004) (Morphology, Experimental, 10) Oleynik, G.S., Shulzhenko, A.A., Stratiychuk, D.A., Gargin, V.G., Vereshchaka, V.M., “Special Features of the Microstructure of a Composite with an Increased Fracture Toughness Produced in the B-C-Si System at High Pressure”, J. Superhard Mater., 26(4), 13–24 (2004), translated from Sverkhtverd. Mater. (Ukraine), 26(4), 16–28 (2004) (Morphology, Experimental, Mechan. Prop., 45) cited from abstract Pai, Ch.-H., “Thermoelectric Properties of Boron Compound-Doped α-SiC Ceramics”, J. Ceram. Soc. Jpn., 112(2), 88–94 (2004) (Crys. Structure, Morphology, Experimental, Electr. Prop., 12) Roger, J., Babizhetskyy, V., Halet, J.-F., Gue´rin, R., “Boron-Silicon Solid Solution: Synthesis and Crystal Structure of a Carbon-Doped Boron-rich SiBn (n 30) Compound”, J. Solid State Chem., 177(11), 4167–4174 (2004) (Crys. Structure, Experimental, 27) Uehara, M., Shiraishi, R., Nogami, A., Enomoto, N., Hojo, J., “SiC-B4C Composites for Synergistic Enhancement of Thermoelectric Property”, J. Eur. Ceram. Soc., 24, 409–412 (2004) (Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., 5) Lee, K.S., Han, I.S., Chung, Y.H., Woo, S.K., Lee, S.W., “Hardness and Wear Resistance of Reaction Bonded SiC-B4C Composite”, Mater. Sci. Forum, 486-487, 245–246 (2005) (Phase Relations, Experimental, Mechan. Prop., 7) Tkachenko, Yu, G., Britun, V.F., Prilutskii, E.V., Yurchenko, D.Z., Bovkun, G.A., “Structure and Properties of B4C-SiC Composites”, Powder Metall. Met. Ceram., 44(3-4), 196–201 (2005) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 4) Fanton, M.A., Cavalero, R.L., Ray, R.G., Weiland, B.E., Everson, W., Snyder, D., Gamble, R., Oslosky, E., “Growth of SiC Boules with Low Boron Concentration”, Mater. Sci. Forum, 527-529, 47–50 (2006) (Morphology, Experimental, Kinetics, 21) Hayun, S., Frage, N., Dariel, M.P., “The Morphology of Ceramic Phases in BxC-SiC-Si Infiltrated Composites”, J. Solid State Chem., 179, 2875–2879 (2006) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, 11) Shulzhenko, A.A., Stratiychuk, D.A., Oleinik, G.S., Vereshchaka, V.M., “Structural Transformations in the Formation of a Superhard Composite in the B-C-Si System”, J. Superhard Mater., 28(4), 11–26 (2006) (Morphology, Phase Relations, Experimental, Interface Phenomena, Mechan. Prop., 26) cited from abstract


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[2007Ber]

[2007Hwa]

[2007Lat]

[2007Mic]

[SGTE] [Mas2] [V-C2]

B–C–Si Ye, Y.-J., Zhang, L.-T., Cheng, L.-F., Xu, Y.-D., “Diffusion Behavior in Amorphous Si-B-C System by Molecular Dynamics Simulation”, J. Inorg. Mater., 21(4), 843–847 (2006) (Morphology, Experimental, Interface Phenomena, Phys. Prop., 13) cited from abstract Berjonneau, J., Langlais, F., Chollon, G., “Understanding the CVD Process of (Si)-B-C Ceramics Through FTIR Spectroscopy Gas Phase Analysis”, Surf. Coat. Technol., 201, 7273–7285 (2007) (Phase Relations, Experimental, Kinetics, 39) Hwang, G.-C., Mastushita, J., Lee, J.-J., “Preparation of Si-B-C System Powder Using Silicon, Boron, and Boron Carbide”, Mater. Sci. Forum, 544-545, 933–936 (2007) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 6) Lattemann, M., Ulrich, S., “Investigation of Structure and Mechanical Properties of Magnetron Sputtered Monolayer and Multilayer Coatings in the Ternary System Si-B-C”, Surf. Coat. Technol., 201(9-11), 5564–5569 (2007) (Morphology, Phase Relations, Experimental, Mechan. Prop., 60) Michaux, A., Sauder, C., Camus, G., Pailler, R., “Young’s Modulus, Thermal Expansion Coefficient and Fracture Behavior of Selected Si-B-C Based Carbides in the 20–1200˚C Temperature Range as Derived from the Behavior of Carbon Fiber Reinforced Microcomposites”, J. Eur. Ceram. Soc., 27(12), 3551–3560 (2007) (Crys. Structure, Calculation, Experimental, Mechan. Prop., Phys. Prop., 21) “Scientific Group Thermodata Europe”, Grenoble Campus, 1001 Avenue Centrale, BP66, F-38402 Saint Martin D’Heres France (Thermodyn., Review) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

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Boron – Carbon – Tantalum Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl

Introduction Although borides and carbides of tantalum are among compounds with the highest melting temperatures (Tm, TaC = 3985˚C, Tm, TaB = 3090˚C), high temperature phase relations in the B-C-Ta system have not yet been explored in detail. Studies on the interaction between tantalum carbides, tantalum borides, B4C and carbon on samples hot-pressed at 1500 < T < 2900˚C [1952Gla, 1955Bre] were used to define phase relations in the B-C-Ta system particularly for the regions TaC-TaB2-B-C and Ta2C-Ta2B [1955Bre]. Further studies of phase equilibria were presented by [1963Rud, 1965Lev, 1976Ord, 1987Ord] and refer to an isothermal section at 1750˚C [1963Rud] and three eutectic quasibinaries: TaB2-C [1965Lev], TaB2-TaC1–x [1976Ord] and TaB2-‘B4C’ [1987Ord]. These results were experimentally accomplished employing X-ray powder diffractometry [1952Gla, 1955Bre, 1963Rud, 1965Lev, 1976Ord, 1987Ord], micrographic analysis [1963Rud, 1965Lev, 1976Ord, 1987Ord], pyrometric melting point measurements [1965Lev, 1976Ord, 1987Ord] and chemical analysis [1963Rud, 1976Ord]. The most relevant data on the topology of the B-C-Ta system were compiled by [1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information up to 1996 was assessed in a general review of phase relations for metal-boron-carbon systems [1998Rog]. Experimental details for all investigations in the B-C-Ta system are summarized in Table 1.

Binary Systems The B-Ta binary system is essentially taken from [1992Rog], however, phase relations in the Ta rich part have been revised recently [2006Cha]: the composition of the eutectic, l Ð (Ta) + Ta2B, has been found at 18 at.% B and the composition of the liquid for the peritectic, l + TaB Ð Ta2B, has been located at 22.5 at.% B. Furthermore the decomposition of Ta2B was found to occur between 1900 and 1950˚C. Due to reinvestigation of C-Ta phase relations from 1700 to 2300˚C by means of diffusion couples, XPD, LOM and light atom EMPA [1996Len, 1997Len, 1998Wie] the C-Ta binary phase diagram of [Mas2] has to be modified. The revised version, as presented in [1998Rog], is shown in Fig. 1. According to calculations employing the order parameter functional method, an ordered NaCl-type derivative phase Ta6C5 is supposed to form below 1150˚C [1991Gus]. A thermodynamic estimation of the systems B-Ta, C-Ta is from [1991Kau]. The B-C system is adopted from an assessment and thermodynamic modelling by [1996Kas, 1998Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The disputed peritectic boron rich reaction L+‘B4C’ Ð (βB)


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was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan].

Solid Phases No ternary compounds were found to exist [1952Gla, 1955Bre, 1963Rud, 1965Lev, 1976Ord, 1987Ord] and mutual solid solubilities at temperatures below 2000˚C were said to be insignificantly low. According to [1963Rud] the maximal solubility of B in the tantalum carbides at 1750˚C was given as less than 1 at.% B. Whereas practically no reaction was observed for TaB2TaC1–x composites up to 2100˚C, a maximal solid solubility of < 2.8 mol% TaB2 in TaC1–x(= 3 mass% TaB2) was reported at 2400˚C raising to 6.7 mol% TaB2 at 2730˚C (= 7 mass% TaB2) [1976Ord] (see also “Quasibinary Sections”). Crystal data for the binary boundary phases are listed in Table 2.

Quasibinary Systems Three quasibinary sections of the eutectic type were reported of which TaB2-TaC0.89 is presented in Fig. 2 (after [1976Ord]) and TaB2-B4.5C is shown in Fig. 3 (after [1987Ord]). Small changes were made to comply with the accepted binary systems. Although the TaC1–x starting powder used by [1976Ord] contained 6.11 mass% C (equivalent to a composition in at.% of Ta50.5C49.5) which is at the carbon rich boundary of the monocarbide rather than at its congruent melting point at 47 at.% C (TaC0.89, Fig. 1), the melting point of TaC1–x was given as 3985˚C [1976Ord] corresponding to the maximum melting point of TaC0.89. From a cursory investigation of the TaB2-C system, [1965Lev] suggested a quasibinary eutectic with the absence of significant mutual solid solubilities. The eutectic point was mentioned at about 32 mol% TaB2 and at 2650˚C.

Invariant Equilibria A tentative partial reaction scheme (see Fig. 4, Table 3) for the B, C rich part was assigned [1998Rog] on the basis of the three experimentally established quasibinary isopleths assuming the formation of three ternary eutectic invariants E1, E2, E3, in correspondence with the phase triangulation in the isothermal section at 1750˚C [1963Rud]; (see also Fig. 5). Temperatures given in Fig. 4 are, however, tentative and strongly depend on the accuracy of the experimentally determined eutectic maxima e2, e3, e5. Depending on the relative temperatures of the reactions p1, e6 and E3, the reaction E3 may alternatively be a transition reaction U1: L + ‘B4C’ Ð TaB2 + (βB) at 2070˚C.

Liquidus, Solidus and Solvus Surfaces No experimentally determined liquidus surface is available. Attempts to schematically represent solidification behavior in the B, C rich part of the system on the basis of the experimentally defined quasibinary reactions e2, e3, e5 (see Fig. 4), cast severe doubts on the correct DOI: 10.1007/978-3-540-88053-0_22 ß Springer 2009

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position of the TaB2-C quasibinary eutectic e3 at 2650˚C [1965Lev], which appears too close to graphite and will need further experimental verification.

Isothermal Sections Figure 5 presents the only isothermal section established for the B-C-Ta ternary system at 1750˚C. The phase triangulation, as obtained by [1963Rud], was amended with additional tie lines to include the phases Ta5B6, Ta3B2 and to comply with the accepted binaries. As the high temperature phase Ta2B was shown by many researchers (see i.e. [2006Cha]) to decompose below 1900˚C it is not included in Fig. 5. As a ternary sample with 5 at.% B near Ta4C3–x did not reveal Ta4C3–x [1963Rud], the likely two-phase region TaB + Ta4C3–x is shown with dashed lines. Alternatively, a very narrow three-phase equilibrium Ta2C + Ta4C3–x + TaC1–xBx may prevent Ta4C3–x to appear at higher boron concentrations. The phase triangulation confirms the observation [1952Gla, 1955Bre] that all tantalum borides with boron contents lower than TaB2 are unstable when heated in combination with carbon.

Thermodynamics No experimental thermodynamic data are presently available for the ternary system. Upper limits for the heat of formation of TaB2 and the lower tantalum borides were estimated from compatibility studies among tantalum borides and carbon [1955Bre].

Notes on Materials Properties and Applications Microhardness of samples near the quasibinary eutectic compositions TaB2 - TaC1–x and TaB2 -‘B4C’ was observed to lie markedly below the linear combination of the two constituents (21.6 GPa for TaB2/TaC1–x and 27 GPa for TaB2/‘B4C’) depending strongly on the dispersion of phases, decreasing with increasing dispersion [1976Ord, 1987Ord]. Engeldinger et al. [1977Eng] demonstrated an increase up to 1300˚C of the hot-hardness of TaC1–x hot-pressed at 2100˚C by adding a few percent of boron. Boron was said to exert a twofold effect: (i) increase of hardness within the solubility range TaC1–xBx (given as 3at.% B maximally) and (ii) hardness increase in the two-phase field (TaC1–x + TaB2) as a combined effect of nonstoichiometry and particle hardening (precipitation of needles and platelets of TaB2 at a coherent interface (100)TaB2//(111)TaC).

Miscellaneous The authors of [1980Ord, 1993Ord] analyzed the interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - B4C [1993Ord] for transition elements M = Ti, Zr, Hf, V, Nb and Ta and presented correlations between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms involved.


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Using a compact diffraction reaction chamber, [2006Won] studied with time resolved X-ray diffraction the chemical dynamics at the combustion front for the directly ignited reaction 3Ta + B4C = 2TaB2 + TaC. Combustion front velocity was 2 mm·s–1 and the adiabatic combustion temperature calculated (2476 K) was well below the Tm of all reactants and reaction products. The combustion (involving no liquid phase) completed in 0.6 s. The authors of [2006Sot] studied the conditions to produce thin B-C-Ta films with variable compositions using a hot Ta-filament (>2000˚C) in a CH4+B2C6+H2 environment: from Auger spectroscopy and XPD the films were said to consist of TaC and Ta2B. Via precipitation strengthening Ta(B,C)-additions increased high-temperature strength and low-temperature ductility in chromium [1975Klo]. Additions of C and/or B4C proved to be an efficient densification aid to obtain up to 98% dense sinter-bodies of tantalum monocarbide (0.36 mass% B4C or 0.43 mass% B4C+0.13 mass % C, hot-pressed at 2200˚C) [2007Zha].

. Table 1 Investigations of the B-C-Ta Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied

Reference


[1952Gla]

Hot-pressing of powder mixtures (powders of TaH2, TaC, TaB and TaB2,) in graphite dies at 1500 to 2900˚C. XPD on nine hot-pressed samples.

Compatibility of TaB2 with C and with ‘B4C’ below 2900˚C. TaB and C or TaC with B or ‘B4C’ always resulted in TaB2 + C.

[1955Bre]

Reaction between metal borides and graphite. Sintering of powder compacts in Mo crucibles under 0.5 bar argon for 50 min at 1777˚C

No experimental data listed but isothermal section shown.

[1963Rud] 45 alloys were prepared by short duration Isothermal section at 1750˚C. hot-pressing in graphite dies at 1200˚C to 2450˚C starting from well blended powders of boron (94 mass% residue O, C, Fe), lampblack C, tantalum (purity 99.7 mass% Ta, containing 0.21 mass% Nb, 0.014% Fe, 0.003% W, 0.036 mass% C) and pre-reacted tantalum carbide (containing 6.34 mass% C of which 0.14% were free C). The powder compacts were heat treated in a W tube vacuum furnace (2.5 Pa) for 9 h at 1750˚C. XPD and LOM.

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[1965Lev]

TaB2 was prepared by vacuum sintering (Ta+2B) powder compacts for 5 h at 1400 to 1700˚C and analyzed by XPD, LOM. Starting materials were Ta powder of 5 to 10 μm containing 1 mass% Nb, 0.93% C, 0.128% Fe, 0.15% Pb, 0.07% Si, 0.07% Ti. Boron impurities were: 0.0036 mass% Fe, 0.0036% Si, 0.0003% Mg, 0.01% Cu, 0.0004% Al, 0.0006% Pb.

Temperature/Composition/Phase Range Studied Interaction of TaB2 with C yielded a quasieutectic system TaB2+C with eutectic point at 32 mol% C and TE of ca 2650˚C. Thermal analysis was performed by direct electrical heating of a graphite tube filled with TaB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.

[1976Ord] Samples were prepared from TaB2 (>89.25 mass% Ta and 10.61% total B and 0.05% free C, 0.02% O2+N2) and TaC (synthesized from 93.73 mass% Ta, 6.11% total C, 0.08% free C, 0.013% O2+N2). Specimens in form of cylinders (3 mm diameter · 50 mm length) were compacted with aid of 15% aqueous polyvinyl alcohol, pre-sintered at 2000˚C in a stream of pure argon prior to heat treatment at above 2400˚C. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

Investigation of the quasieutectic system TaB2+TaC1–x with eutectic point at 66 mass% TaB2 and TE of 2730±40˚C on 13 samples.

[1987Ord] Samples were prepared from TaB2 and B4C which was vacuum annealed at 2000˚C to reduce C-content to 0.2 mass% free C. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

Investigation of the quasieutectic system TaB2+‘B4C’ with eutectic point at 31–33 mol% TaB2 and TE of 2370±30˚C on 12 samples.


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[2005Tan] Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)); the liquidus and solidus curves to the L+(βB) field have been derived via chemical analysis.

Temperature/Composition/Phase Range Studied Confirmation of peritectic type of reaction L+B4+xC Ð (βB) via determination of the liquidus and solidus curves to the L+(βB) field.

. Table 2 Crystallographic Data of Solid Phases Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype


Comments/References

(Ta) < 3020

cI2 Im 3m W

a = 330.3

at 25˚C [Mas2]

(βB) < 2092

hR333 R 3m βB

a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1092.05 c = 2386.73

[1993Wer]

hP4 P63/mmc C (graphite)

a = 246.12 c = 670.90 a = 246.023 c = 671.163 a = 246.75 c = 669.78

(C)gr < 3827 (S.P.)

at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x at TaB99.5 [1992Rog] at 25˚C [Mas2] [1967Low] at 2.35 at.% Cmax (2350˚C) linear ∂a/∂x, ∂c/∂x, [1967Low]

(C)d

cF8 Fd 3m C (diamond)

a = 356.69

at 25˚C, 60 GPa [Mas2]

‘B4C’ < 2450

hR45 R 3m B13C2

a = 565.1 c = 1219.6 a = 560.7 c = 1209.5 a = 556.0 c = 1212.0 a = 556.1 c = 1212.0

9 to 20 at.% C [1990Ase]

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quenched from 2400˚C [1987Ord] sample B4C+87 mass% TaB2 quenched from 2430˚C [1987Ord]


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. Table 2 (continued) Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype B25C

tP68 P 42m B25C or tP68 P42/nnm B25C


Comments/References

a = 875.3 ± 0.4 c = 509.3 ± 1.5


a = 875.3 ± 0.4 c = 509.3 ± 1.5

[V-C2] defect structure tP52 [1998Rog]

Ta2B 2417 - 1900 [2006Cha]

tI12 I4/mmm CuAl2

a = 577.93 c = 486.38

[1992Rog]

Ta3B2 < 2180

tP10 P4/mbm U3Si2

a = 619.27 c = 330.27

[1992Rog]

TaB < 3090

oC8 Cmcm CrB

a = 327.49 b = 868.16 c = 315.84

[1992Rog]

Ta5B6

oC22 Cmmm V5B6

a = 313.85 b = 2260.2 c = 328.95

[1992Rog]a)

Ta3B4 < 3030

oI14 Immm Ta3B4

a = 313.20 b = 1399.68 c = 328.84

[1992Rog]a)

TaB2 < 3037

hP3 P6/mmm AlB2


a = 309.73 c = 322.57 a = 305.86 c = 328.92 a = 309.8 c = 322.5 a = 310.0 c = 323.0 a = 311.0 c = 325.0 a = 312.5 c = 326.0 a = 307.8 c = 326.5 a = 307.9 c = 326.4 a = 308.0 c = 326.4

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26 to 37 at.% B [1992Rog] Ta rich [1992Rog] B rich [1992Rog] at 25˚C (298 K) [V-C2] at 227˚C (500 K) [V-C2] at 727˚C (1000 K) [V-C2] at 1227˚C (1500 K) [V-C2] sample quenched from 3100˚C [1976Ord] sample TaB2+5 mass% TaC1–x quenched from 2900˚C [1976Ord] sample TaB2+7.6 mass% B4C quenched from 3000˚C [1987Ord]

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. Table 2 (continued) Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype Ta2C(h) 3330 - 2020

Ta2C(r) ≤ 2020

Ta4C3–x ≤ 2170

hP4 P63/mmc defect NiAs

hP3 P 3m1 CdI2

hR24 R 3m V4C3–x

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a = 310.5 c = 494.0 a = 310.5 c = 494.5

Comments/References 26 to 35.6 at.% C [1996Len, 1997Len, 1998Wie] defect structure hP3 [1998Rog] at Ta2C0.92 [V-C2] quenched from 1750˚C [1963Rud]

a = 310.37 ± 0.04 at 25˚C (298 K) [V-C2] c = 493.94 ± 0.11 a = 310.5 c = 494.1 a = 311.4 c = 495.3 a = 312.6 c = 496.8

a = 311.6 ± 0.5 c = 3000 ± 5

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at 227˚C (500 K) [V-C2] at 627˚C (900 K) [V-C2] at 1127˚C (1400 K) [V-C2] 38.2 to 39.0 at.% C [1996Len, 1997Len, 1998Wie] defect structure hR20 [1998Rog] [V-C2]


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. Table 2 (continued) Pearson Phase/ Symbol/ Temperature Range Space Group/ [˚C] Prototype TaC1–xb) < 3985

cF8 Fm 3m NaCl


Comments/References

36.5 to 49.8 at.% C [1996Len, 1997Len,1998Wie] TaC0.74 to TaC1.0, a = 441.3 to 445.4 quenched from 1750˚C [1963Rud] a = 442.43 at TaC0.789, 298 K [V-C2] a = 445.62 ± 0.02 at TaC0.997, 298 K [V-C2] a = 446.07 at TaC0.997, 474 K [V-C2] a = 446.67 at TaC0.997, 696 K [V-C2] a = 447.26 at TaC0.997, 891 K [V-C2] a = 447.95 at TaC0.997, 1087 K [V-C2] a = 445.1 quenched from 3980˚C [1976Ord] a = 445.32 sample TaC1–x+2.85 mol% TaB2 quenched from 3650˚C [1976Ord] a = 445.40 sample TaC1–x+4.75 mol% TaB2 quenched from 3620˚C [1976Ord] a = 445.57 sample TaC1–x+6.68 mol% TaB2 quenched from 3590˚C [1976Ord] a = 445.57 sample TaC1–x + 9.54 mol% TaB2 quenched from 3400˚C [1976Ord]

a)

Note: Crystal setting standardized with program Typix [1994Par]. Note: The symmetry assigned to the (hypothetical) structure Ta6C5 was C2, C2/m or P31 deriving from the parent NaCl type phase TaC1–x as a superlattice structure [1991Gus].

b)

. Table 3 Invariant Equilibria Composition (at.%) Reaction L Ð TaB2 + TaC1–x

T [˚C]

Type

2730±40

Phase

B

C

Ta

e2(max)

L

49.8

11.9

38.3

L Ð TaB2 + (C)gr

2650

e3(max)

L

40

40

20

L Ð TaB2 + TaC1–x + (C)gr

2550

E1

-

-

-

-

L Ð TaB2 +’B4C’

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e5(max)

L

77.5

14.5

8

L Ð TaB2 +’B4C’ + (C)gr

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E2

-

-

-

-

L Ð TaB2 +’B4C’ + (βB)

2000

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-

-

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. Fig. 1 B-C-Ta. The C-Ta phase diagram

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. Fig. 2 B-C-Ta. Vertical section TaB2-TaC1–x


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. Fig. 3 B-C-Ta. Vertical section TaB2 -‘B4C’

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. Fig. 4 B-C-Ta. Partial reaction scheme (proposed)

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. Fig. 5 B-C-Ta. Isothermal section at 1750˚C

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References [1952Gla] [1955Bre] [1963Rud]

[1965Lev]

[1967Low] [1975Klo] [1976Ord]

[1977Eng]

[1980Ord]

[1983Sch]

[1984Hol]

[1987Ord]

[1990Ase]

[1991Gus]

[1991Kau] [1992Rog]

[1993Ord] [1993Wer]

Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–406 (1955) (Experimental, Thermodyn., 19) Rudy, E., Benesovsky, F., Toth, L.E., “Investigation of Ternary System Between Va and VIa-Metals with Boron and Carbon” (in German), Z. Metallkd., 54(6), 345–353 (1963) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, 43) Levinskii, Y.V., Salibekov, S.E., Levinskaya, M.K., “Interaction of Diborides of V, Nb, Ta with Carbon” (in Russian), Poroshk. Metall. (Kiev), 5(11), 66–69 (1965) (Experimental, Phase Diagram, Phase Relations, 6) Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–144 (1967) (Crys. Structure, Experimental, 5) Klopp, W.D., “A Review of Chromium, Molybdenum, and Tungsten Alloys”, J. Less-Common Met., 42(3) 261–278 (1975) (Experimental, Mechan. Prop., 51) Ordanyan, S.S., Unrod, V.I., Polishchuk, V.S., Storonkina, N.M., “Reactions in the System TaC-TaB2”, Powder Metall. Met. Ceram., 15(9), 692–695 (1976), translated from Poroshk. Metall., 9(165), 40–43 (1976) (Crys. Structure, Kinetics, Morphology, Phase Relations, Phase Diagram, #, 6) Engeldinger, M., Ritzhaupt-Kleissl, H.J., Thuemmler, F., “Hard Materials in the Ta-C-B System”, Sci. Sintering, 9(1), 121–140 (1977) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, Phys. Prop., 29) Ordanyan, S.S., “Laws of Interaction in the Systems MIV, VC-MIV, VB2”, Inorg. Mater., 16(8), 961–965 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(8), 1407–1411 (1980) (Experimental, Thermodyn., 14) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron Systems” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 264–274 (Review, Crys. Structure, Phase Diagram, Phase Relations, 87) Ordanyan, S.S., Dmitriev A.I., Bizhev, K.T., Stepanenko, E.K., “The Interaction in B4C-MeVB2 Systems”, Powder Metall. Met. Ceram., 26(10), 834–836 (1987), translated from Poroshk. Metall., 10(298), 66–69 (1987) (Morphology, Phase Diagram, Phase Relations, #, 5) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Gusev, A.I., “Phase Diagrams for Ordering Systems in the Order-Parameter Functional Method”, Sov. Phys. Solid State, 32(9), 1595–1599 (1991) (Theory, Phase Diagram, Phase Relations, Thermodyn., 18) see also Gusev, A.I., “Physical Chemistry of Nonstoichiometric Refractory Compounds” (in Russian), Chapter 3, Nauka, Moscow, (1991) (Review, Thermodyn., Crys. Structure, Phase Diagram, Phase Relations, 102) Kaufman, L., “Coupled Thermochemical and Phase Diagram Data for Tantalum Based Binary Alloys”, Calphad, 15(3), 243–259 (1991) (Phase Diagram, Phase Relations, Thermodyn., 27) Rogl, P., “The System B-N-Ta” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems”, Rogl, P., Schuster, J.C. (Eds.), ASM International, Materials Park, Ohio, USA, 97–100 (1992) (Crys. Structure, Thermodyn., Phase Diagram, Phase Relations, Experimental, Review, 7) Ordan’yan, S.S., “On Regularities of Interaction in the Systems B4C - MeIV - MeVIB2” (in Russian), Ogneupory, 1, 15–17 (1993) (Phase Diagram, Phase Relations, Review, Theory, 18) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51)


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[1994Par]

[1995Vil] [1996Len]

[1996Kas] [1997Len]

[1998Kas]

[1998Rog]

[1998Wie]

[2005Tan]

[2006Cha]

[2006Sot]

[2006Won]

[2007Zha]

[Mas2] [V-C2]

B–C–Ta McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 190–191 (1994) (Phase Diagram, Phase Relations, Review, 2) Parthe, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1–4, Gmelin, Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA, 5366–5370 (1995) (Review, Phase Diagram, Crys. Structure, 7) Lengauer, W., Wiesenberger, H., Joguet, M., Rafaja, D., Ettmayer, P., “Chemical Diffusion in Transition Metal - Nitrogen Systems” in “The Chemistry of Transition Metal Carbides and Nitrides”, Oyama, S.T. (Ed.), Blacky Academic, Oxford, 91–106 (1996) (Phase Diagram, Phase Relations, Experimental, #, 29) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Lengauer, W., Wiesenberger, H., Mayr, W., Bidaud, E., Berger, R., Ettmayer, P., “Phase Stabilities of Transition Metal Carbides and Nitrides Investigated by Reaction Diffusion”, J. Chim. Phys., 94, 1020–1025 (1997) (Experimental, Phase Relations, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Tantalum” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM International, Materials Park, Ohio, USA, 257–268 (1998) (Experimental, Crys. Structure, Review, Phase Diagram, Phase Relations, 22) Wiesenberger, H., Lengauer, W., Ettmayer, P., “Reaction Diffusion and Phase Equilibria in the V-C, Nb-C, Ta-C and Ta-N Systems”, Acta Mater., 46(2), 451–666 (1998) (Experimental, Phase Diagram, Phase Relations, 30) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC⊊Ð⊊βB”, research presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg, Germany, August 21–26, 142 (2005) (Experimental, Phase Diagram, Phase Relations, 4) Chad, V.M., Ramos, E.C.T., Coelho, G.C., Nunes, C.A., Suzuki, P.A., Flavio, F., Rogl, P., “Evaluation of the Invariant Reactions in the Ta-rich Region of the Ta-B System”, J. Phase Equilib. Diffus., 27(5), 452–455 (2006) (Experimental, Phase Diagram, Phase Relations, Crys. Structure, 10) Soto, G., Silva, G., Contreras, O., “A Study on the Flexibility of the Hot-Filament Configuration and its Implementation for Diamond, Boron Carbide and Ternary Alloys Deposition”, Surf. Coat. Technol., 201 (6), 2733–2740 (2006) (Experimental, Morphology, Phase Relations, 22) Wong, J., Larson, E.M., Waide, P.A., Frahm, R., “Combustion Front Dynamics in the Combustion Synthesis of Refractory Metal Carbides and di-Borides Using Time-Resolved X-ray Diffraction”, J. Synchrotron Radiat., 13(pt.4), 326–335 (2006) (Experimental, Phys. Prop., 30) Zhang, X.H., Hilmas, G.E., Fahrenholtz, W.G., Deason, D.M., “Hot Pressing of Tantalum Carbide with and without Sintering Additives”, J. Am. Ceram. Soc., 90(2), 393–401 (2007) (Experimental, Mechan. Prop., 24) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

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Boron – Carbon – Vanadium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl

Introduction Interaction of vanadium carbides and borides is not only of significance for the design of borided vanadium rich steels but also for superhard boride-carbide composites. Although several research groups have studied phase relations in the B-C-V system [1952Gla, 1963Rud, 1965Lev, 1975Mar, 1982Ord, 1987Ord], hitherto neither high temperature data are available nor a liquidus projection for the complete B-C-V ternary system. Based on a preliminary evaluation [1952Gla] of the chemical interactions among vanadium borides (VB, VB2), carbon and boron carbide in samples hot-pressed at T >1500˚C, an isothermal section at 1450˚C [1963Rud] and three quasibinary eutectic sections: VB2-C [1965Lev], VB2-VC1–x [1982Ord] and VB2-‘B4C’ [1987Ord], were established experimentally. Techniques of analysis employed were: X-ray powder diffractometry [1952Gla, 1963Rud, 1965Lev, 1975Mar, 1982Ord, 1987Ord], micrographic inspection [1963Rud, 1965Lev, 1982Ord, 1987Ord], melting point analyses [1965Lev, 1982Ord, 1987Ord] and chemical analyses [1963Rud, 1982Ord]. The most relevant data on the topology of the B-C-V system were compiled by [1983Sch, 1984Hol, 1994McH, 1995Vil]. A full status of all information up to 1996 was assessed in a general review of phase relations for metal-boron-carbon systems [1998Rog] including a thermodynamic extrapolation of the B-C-V ternary system on the basis of binary data known at that time. Some of the calculated sections were reproduced in [1999Rog]. Experimental details for all investigations in the B-C-V system are summarized in Table 1.

Binary Systems The C-V binary equilibria still need reinvestigation particularly with respect to the melting point of VC1–x as well as with respect to the behavior in the region V2C-VC1–x at temperatures below about 1200˚C. Detailed investigation of the C-V phase relations from 1200 to 1900˚C by means of diffusion couples, XPD, LOM and light atom EMPA [1996Len, 1997Len, 1998Wie] invokes a revision of the C-V binary given by [Mas2, 1985Car]. Furthermore, the C-V diagram presented in [Mas2, 1985Car] is essentially based on the experimental data by [1969Rud], although the melting point of vanadium monocarbide was raised from 2648 ± 12˚C at 43.0 ± 0.5 at.% C [1969Rud] to 2800˚C at 45.5 at.% C [Mas2, 1985Car]. Accordingly the eutectic l Ð VC1–x + (C)gr was raised from 2625 ± 12˚C at 49.5 ± 0.5 at.% C [1969Rud] to 2670˚C at 53.5 at.% C [Mas2, 1985Car]. A thermodynamic assessment and calculation of the C-V system by [1991Hua] arrived at 2656˚C and 44.62 at.% C for the melting point of VC1–x and at 2605˚C, 50.6 at.% C for the VC1–x + (C)gr eutectic (see Fig. 1a). The authors of [1991Gus] calculated the formation of ordered NaCl type derivative phases Landolt‐Bo¨rnstein New Series IV/11E1

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V3C2, V6C5 and V8C7 below ca. 1200˚C employing the order parameter functional method. Besides the diffusion data of [1998Wie], phase equilibria studies by XPD on hot-pressed and annealed alloys [1999Lip] were said to yield a peritectoid formation V2C+VC1–xÐ V6C5 at 1212˚C, a peritectoid formation V2C+V6C5ÐV3C2 at 882˚C and a V8C7 superstructure of VC1–x at the C rich boundary below 1107˚C. However, no clear decision was taken, how the V4C3–x phase enters the phase relations; a decomposition of V4C3–x below 1200˚C was suggested [1999Lip], but carbon contents of the ordering phases seem to be too high with respect to the low-carbon phase boundary of the monocarbide as established from diffusion couples [1998Wie]. A tentative revised version of the C-V system, which tries to interrelate all data hitherto reported is shown in Fig. 1b. A thermodynamic calculation of the B-V system is from [1981Spe] with a recent refinement of this modeling by [1998Rog, 2001Fab]. All these calculations reproduced the peritectoid formation V3B4+VB Ð V5B6 at about 1727˚C, as given by [1969Spe]. A high temperature study, however, revealed V5B6 forming from the liquid in a peritectic reaction l + V3B4 Ð V5B6 at about 2600˚C [2004Nun]. Consequently the formation of V5B6 from the liquid implies a change in the peritectic formation of VB: l + V5B6 Ð VB (see Fig. 2). The B-C system is adopted from an assessment and thermodynamic modeling by [1996Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The disputed peritectic boron rich reaction L+‘B4C’⊊Ð (βB) was experimentally confirmed from a floating zone experiment on several carbon-doped boron samples [2005Tan]. The crystallographic data of all phases pertinent to the phase equilibria are listed in Table 2.

Solid Phases Besides the very boron rich V0.65B24C [1980Amb] no ternary B-C-V compound was reported [1952Gla, 1963Rud, 1965Lev, 1975Mar, 1982Ord, 1987Ord]. Furthermore, mutual solid solubilities among vanadium borides and carbides were observed to be very low. The only exceptions are VC1–x and V3B2. For V3B2 a carbon solubility of 3 at.% C at 1450˚C was estimated by [1963Rud] from the significantly lower lattice parameters in the ternary system and accordingly a B/C substitutional solution, V3(B1–xCx)2, was assumed. The solubility of B in VC1–x at 1450˚C was observed not to exceed about 1 at.% B [1963Rud]. Samples from the quasibinary system VB2+VC0.8, however, revealed an extended solid solubility of about 10 mass% VB2 in congruently melting VC1–x at the temperature of the quasibinary eutectic 2120 ±20˚C [1982Ord]. The solubility was said to quickly drop to 5 mass% VB2 at 2000˚C, and at 1800˚C unit cell dimensions of VC0.88 remain unchanged [1982Ord]. Although no variation of lattice parameters was encountered on B4C in B4C-VB2 samples [1987Ord], a slight increase of the unit cell dimensions of B4C alloyed with V was said to be indicative of a small solubility of V in B4C [1975Mar].

Quasibinary Systems Three quasibinary sections of the eutectic type were experimentally established of which two are presented in Fig. 3 (VB2-VC0.862, after [1982Ord]) and Fig. 4 (VB2-‘B4C’, after [1987Ord]), with small changes to comply with the accepted binary systems. With respect to the rather low DOI: 10.1007/978-3-540-88053-0_23 ß Springer 2009

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melting points of the boundary compounds recorded by [1982Ord] (VB2, 2700˚C; VC1–x, 2650˚C) in comparison with the accepted values (VB2, 2750˚C; VC0.85, 2800˚C) (Figs. 1a, 1b and 2), the quasibinary eutectic temperature TE = 2120±20˚C [1982Ord] seems low. Furthermore, VC0.88 was claimed to be the composition exhibiting congruent melting behavior [1982Ord], whilst [Mas2] actually assessed VC0.85. Preliminary data on the VB2-C system state the quasibinary eutectic nature of the isopleth (see Fig. 5) with insignificant mutual solubilities [1965Lev]. The eutectic point was reported at about 30 mol% VB2 and 2450˚C [1965Lev]. A reinvestigation is recommended.

Invariant Equilibria No experimentally established reaction scheme exists for the entire B-C-V ternary system, nevertheless a tentative partial reaction scheme for the vanadium-poor part can be assigned on the basis of the three experimentally observed ‘eutectic’ quasibinary sections assuming the formation of three invariants (E1, E2, U3, see Figs. 6a and 6b) in correspondence with the phase triangulation at 1450˚C [1963Rud]. Figures 6a and 6b present the reaction scheme as a results of a thermodynamic extrapolation of the entire system [1998Rog]. This reaction scheme has been amended taking care of the higher stability of the V5B6 phase (l + V3B4 Ð⊊V5B6 at about 2600˚C [2004Nun]) and assuming that the field of primary crystallization will run parallel to the primary field of V3B4 until V5B6 finally will be consumed in a ternary eutectic reaction L Ð⊊V5B6 + VB + VC1–x. As in an alternative case the L+V5B6 field may peter out before joining the VC1–x field, for instance in a transition reaction L + V5B6 Ð⊊VB + V3B4, a detailed experimental determination of the ternary reactions and of the liquidus surface is recommended.

Liquidus, Solidus and Solvus Surfaces No liquidus surface was derived experimentally. Experimental observations only concern the three eutectic quasibinary maxima, e3, e4, e5 (see Table 3). From a thermodynamic extrapolation of the binaries the liquidus surface was calculated [1998Rog], however, employing the older B-V binary system with V5B6 only stable below 1727˚C: the liquidus surface projection, as calculated but with a narrow primary field of V5B6 inserted (see discussion in section “Invariant Equilibria”), is shown in Fig. 7. Comparison with the experiments revealed doubts on the correct position and temperature of the quasibinary eutectic reaction e4(max) (L Ð⊊VB2C + (C)gr) [1965Lev]; further experimental verification is recommended.

Isothermal Sections Figure 8 is a presentation of the phase relations calculated at 1450˚C [1998Rog], which compare well with the results of [1963Rud] with some additional tie lines that incorporate the novel binary vanadium borides V5B6, V2B3 in consistency with the accepted binary boundary systems. The phase triangulation confirms the observations of [1952Gla] that all vanadium borides with boron contents lower than VB2 are unstable, when heated in combination with carbon. A series of isothermal sections was calculated thermodynamically Landolt‐Bo¨rnstein New Series IV/11E1

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[1998Rog] of which those at 1600 and 2000˚C are shown in Fig. 9 and Fig. 10, respectively. Tie lines to V5B6 were inserted in the latter to comply with the accepted B-V binary.

Temperature – Composition Sections Several isopleths were calculated [1998Rog, 1999Rog], however, employing the older B-V binary system with V5B6 only stable below 1727˚C. The isopleths VC1–x-B4C and VB2-VC1–x (x = 0.082, carbon rich boundary of VC at low temperature) are not affected by V5B6 and are shown in Figs. 11 and 12.

Thermodynamics A thermodynamic calculation of the ternary B-C-V system [1998Rog, 1999Rog] was based on thermodynamic assessments of the binary systems B-C [1996Kas], B-V [1998Rog] and C-V [1991Hua] as well as relying on the phase diagram data from [1963Rud] for the optimization of the thermodynamic parameters. The solid solubilities of both B and C in (V) could be satisfactorily reproduced by ternary ideal mixing terms in its Gibbs energy description. More accurate experimental investigations are needed to clarify the established shortcomings and to refine the current description of the B-C-V system.

Notes on Materials Properties and Applications [1968Alp] discussed phase assemblages, microstructures, and properties of fused carbides and/or borides from refractory systems (also for B-C-V) containing free graphite in terms of compatibility and phase diagrams. All these bodies have excellent thermal-shock resistance. Other properties (such as electrical, thermal, mechanical, chemical) can be modified by choosing different phase assemblages. Some of these materials have been cast into large shapes more than 45 cm long, and they can be machined into articles. The authors of [1998Rad, 2002Rad] fabricated a new class of superhard boron carbidebased materials by pressureless sintering of B4C-VB2 and/or B4C-VC compacts showing a considerable increase in microhardness (76 GPa) and abrasive wear resistance values of the sintered materials (1.8 times as compared to “pure” hot-pressed B4C). Free carbon and vanadium boride in hot-pressed B4C - VB2 - C composites were reported to activate the sintering process and to obtain dense, highly dispersed ceramics with higher hardness and bending strength than the monophase boron carbide ceramic. The new composite material was said to be promising for fabricating wear-resistant and shock-resistant components of various structures and machines [2006Gri].

Miscellaneous The authors of [1980Ord, 1993Ord] analyzed the interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - B4C [1993Ord] for transition elements M = Ti, DOI: 10.1007/978-3-540-88053-0_23 ß Springer 2009

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Zr, Hf, V, Nb, Ta and presented correlations between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms involved. Microhardness values of samples near the quasibinary eutectic compositions VB2 - VC1–x and VB2 -‘B4C’ were observed to lie markedly below the linear combination of the two constituents (11 to 18 GPa for VB2/VC0.88 and 21.4 to 25.2 GPa for VB2/’B4C’) depending strongly on the dispersion of phases, decreasing with increasing dispersion [1982Ord, 1987Ord]. Physical properties of ‘B4C’ - V, VB2 + ‘B4C’ + B and VB + ‘B4C’ + (C)gr cermets were studied: microhardness, microfracture [1975Mar, 1976Mar], thermal expansion [1975Mar, 1976Mar] and electrical resistance [1975Mar].

. Table 1 Investigations of the B-C-Ta Phase Relations, Structures and Thermodynamics Reference [1952Gla]

Method/Experimental Technique Hot-pressing of powder mixtures (powders of B, C, V, VC, VB and VB2,) in graphite dies at 1500 to 3150˚C. XPD on ten hot-pressed samples.

Temperature/Composition/Phase Range Studied Compatibility of VB2 with C and with ‘B4C’ below 2100˚C. VB and C or VC with B or B4C resulted in VB2 + C.

[1963Rud] 30 alloys were prepared by short duration Isothermal section at 1450˚C. hot-pressing in graphite dies at 1200˚C to 2450˚C starting from well blended powders of boron (94 mass% residue O, C, Fe), lampblack C, and vanadium (purity 99.72 mass% V, containing 0.18 mass% Fe, 0.026% N, 0.13% O, 0.055% C and 0.0053 mass% H). The powder compacts were heat treated in a W-tube vacuum furnace (2.5 Pa) for 9 h at 1450˚C. Alloys near VB2 were annealed for 150 min at 1900˚C. XPD and LOM.


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. Table 1 (continued) Reference [1965Lev]



VB2 was prepared by vacuum sintering Interaction of VB2 with C yielded a of powder compacts for 5 h at 1400 to quasieutectic system VB2+C with eutectic 1700˚C and analyzed by XPD, LOM. point at 30 mol% C and TE of ca. 2450˚C. Starting materials were VH0.7 powder of 5 to 10 μm containing 0.6 mass% Fe, 0.08% Ti, 0.15% Pb, 0.05% Si, 0.16% O. Boron impurities were: 0.0036% Fe, 0.0036% Si, 0.0003% Mg, 0.01% Cu, 0.0004% Al, 0.0006% Pb. Thermal analysis was performed by direct electrical heating of a graphite tube filled with VB2-powder (outer diameter 8 mm, inner diameter 2 mm, length 80 mm) in vacuum to a specified temperature, holding for some time and switching off the furnace. Melting temperature was measured with an optical pyrometer and was taken as the lowest temperature hold after which a frozen drop was observed in the bore of the broken tube.

[1975Mar] Specimens of V+‘B4C’ with 2–5% porosity V+‘B4C’ were prepared by hot pressing at 1900˚C under Ar for 5–7 min and were then stress relieved at 1600˚C for 12 h in vacuum. Phase analyses by XPD [1982Ord] Samples were prepared by intimate Investigation of the quasieutectic system mixing under ethanol of powders of VC0.88 VB2+VC0.88 with eutectic point at 46 mol% VB2 and TE of 2120 ±⊊20˚C on 13 samples. (synthesized from the elements) and commercial VB2 powders. Specimens (3 mm · 3 mm · 40 mm length) were obtained by spark erosion from blanks pre-sintered at 1800˚C for 2 h in vacuum (0.013 Pa). 1 mass% of Ni was added as a densification aid and was said to have evaporated almost completely during high vacuum sintering. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

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Investigation of the quasieutectic system VB2+‘B4C’ with eutectic point at 45–48 mol% VB2 and TE of 2170±30˚C on 13 samples.

[1987Ord] Samples were prepared from commercial powders of VB2 and B4C which were vacuum annealed at 10 mPa and 2000˚C to reduce the C-content to 0.2 mass% free C. Specimens with high B4C content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

Stability of V5B6 forming from the liquid [2004Nun] Investigation of three alloys V48B52, V45.45B54.55, V44B56, prepared by Argon arc l + V3B4 Ð⊊V5B6 melting from the elements (99.75 mass% V, 99.5 mass% B). A part of the samples was vacuum annealed at 2000˚C for 2 h. SEM, XPD. [2005Tan] Floating zone refinement on six rods with starting compositions (B + 0.1 (0.2, 0.3, 1.0, 1.1 and 1.3 at.% C)); the liquidus and solidus curves to the L+(βB) field have been derived via chemical analysis.

Confirmation of peritectic type of reaction L + B4+xC⊊Ð⊊(βB) via determination of the liquidus and solidus curves to the L+(βB) field.


Phase/ Temperature Range [˚C] (V) < 1910


Pearson Symbol/ Space Group/ Prototype cI2 Im3m W

Lattice Parameters [pm] a = 302.40

Comments/References at 25˚C [Mas2]

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Phase/ Temperature Range [˚C] (βB) < 2092

(C)gr < 3827 (S.P.)

Pearson Symbol/ Space Group/ Prototype hR333 R3m βB

Lattice Parameters [pm] a = 1093.30 c = 2382.52 a = 1092.2 c = 2381.1 a = 1094.76 ± 0.07 c = 2384.22 ± 0.23 a = 1097.2 ± 0.3 c = 2390.8 ± 0.9 a = 1094.9 ± 0.3 c = 2384.0 ± 1.0

hP4 a = 246.12 c = 670.90 P63/mmc C (graphite) a = 246.023 c = 671.163 a = 246.75 c = 669.78

(C)d

cF8 a = 356.69 Fd3m C (diamond)

‘B4C’ < 2450

hR45 R3m B13C2

Comments/References [1993Wer] at 1.1 at.% C [1993Wer] linear ∂a/∂x, ∂c/∂x at VB20 [V-C2] at VB65 [V-C2] at VB99 [V-C2] at 25˚C [Mas2] [1967Low] at 2.35 at.% Cmax (2350˚C), linear ∂a/∂x, ∂c/∂x, [1967Low] at 25˚C, 60 GPa [Mas2]

a = 565.1 to 560.7 9 to 20 at.% C [1990Ase] c = 1219.6 to 1209.5 a = 556.0 c = 1212.0 a = 556.1 c = 1212.0

sample quenched from 2400˚C [1987Ord] for 15.5 mol% VB2, 84.5% B4C, quenched from 2420˚C [1987Ord]

B25C

tP68 P42m or P42/nnm B25C

a = 875.3 ± 0.4 c = 509.3 ± 1.5


V3B2 < 1913

tP10 P4/mbm U3Si2

a = 574.14 c = 302.95

[1998Rog]

a = 572.8 to 573.9 c = 302.6 to 303.0 a = 570.3 c = 302.5

[1963Rud]

V3(B1-xCx)2

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sample quenched from 1450˚C [1963Rud] contains also V2C


B–C–V

23





Comments/References

VB < 2549

oC8 Cmcm CrB

a = 306.03 b = 805.4 c = 297.2

[1998Rog]

V5B6 ≲ 2560

oC22 Cmmm V5B6

a = 297.64 b = 2130.8 c = 305.8

[1998Rog, 2004Nun]

V3B4 < 2600

oI14 Immm Ta3B4

a = 298.1 b = 1322.0 c = 306.06

[1998Rog]

V2B3 < 2610

oC20 Cmcm V2B3

a = 305.88 b = 1842.2 c = 298.46

[1998Rog]

VB2 < 2750 (< 2700±50 [1982Ord])

hP3 P6/mmm AlB2

a = 299.89 c = 305.8 a = 299.8 c = 306.0 a = 300.2 c = 306.2 a = 299.7 c = 306.1 a = 300.2 c = 306.5

[1998Rog]

V1-xB25

tP52 P42/nnm TiB25

a = 882.4 ± 0.9 c = 507.2 ± 1.2

[V-C2] metastable (?)

βV2C(h1) < 2187

hP4 P63/mmc NiAs

a = 288.78 c = 457.43 a = 288.0 to 289.4 c = 445.59 to 459.0

27 to 35.4 at.% C [1996Len] defect NiAs structure hP3 [1998Rog] [V-C2] labelled as εFe2N type sample quenched from 1450˚C [1963Rud]

β’V2C(h2) 1600 - 800

hP9 P31ma) W2C

αV2C < 850

oP12 Pbcn Fe2N


sample quenched from 2600˚C [1987Ord] sample quenched from 2520˚C [1982Ord] sample VB2-6.7 mol% B4C quenched from 2400˚C [1987Ord] sample VB2-5 mol% VC0.88, quenched from 2320˚C [1982Ord]

a = 500.5 c = 455.1

[1998Rog]

0

0 -

84.7 15.3

66.8

98.7

1912

9.7 0 1.3

0.1 33.2

1244

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. Fig. 1 B-C-W. Calculated phase diagram B-W

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. Fig. 2 B-C-W. Calculated phase diagram C-W


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. Fig. 3a B-C-W. Reaction scheme, part 1

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. Fig. 3b B-C-W. Reaction scheme, part 2

B–C–W


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. Fig. 3c B-C-W. Reaction scheme, part 3

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. Fig. 4 B-C-W. Liquidus surface projection


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. Fig. 5 B-C-W. Isothermal section at 1500˚C with alloy data from [1970Rud]

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. Fig. 6 B-C-W. Phase relations at subsolidus temperatures with alloy data from [1970Rud]


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. Fig. 7 B-C-W. Calculated isothermal section at 2150˚C

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. Fig. 8 B-C-W. Calculated isothermal section at 2350˚C


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. Fig. 9 B-C-W. Calculated isopleth W2B5–x -‘B4C’ with alloy data from [1970Rud]

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. Fig. 10 B-C-W. Calculated isopleth W2B5–x -C with alloy data from [1970Rud]


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. Fig. 11 B-C-W. Calculated isopleth WB-‘B4C’

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. Fig. 12 B-C-W. Calculated isopleth W-‘B4C’


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References [1952Gla] [1955Bre] [1959Sam]

[1963Rud]

[1965Lev]

[1965Rud]

[1966Kip] [1967Low] [1968Kip1]

[1968Kip2]

[1969Rud1]

[1969Rud2] [1970Rud]

[1973Kuh]

[1978Har] [1983Sch]

[1984Hol]

[1986Gus] [1986Loe]

Glaser, F.W., “Contribution to the Metal-Carbon-Boron System”, J. Metals, 4(4), 391–396 (1952) (Crys. Structure, Experimental, 19) Brewer, L., Haraldsen, H., “The Thermodynamic Stability of Refractory Borides”, J. Electrochem. Soc., 102, 399–406 (1955) (Experimental, Thermodyn., 19) Samsonov, G.V., “The Interaction of Ti, Zr and W Borides with their Carbides”, Vopr. Poroshk. Metall. i Prochnosti Materialov, Akad. Nauk Ukr. SSR, (7), 72–98 (1959) (Experimental, Phase Diagram, Phase Relations, 15) Rudy, E., Benesovsky, F., Toth, L.E., “Investigation of Ternary System Between Va and VIa-Metals with Boron and Carbon” (in German), Z. Metallkd., 54(6), 345–353 (1963) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, #, 43) Levinskii, Yu.V., Salibekov, S.E., Levinskaya, M.K., “Reaction of Chromium, Molybdenum and Tungsten Borides with Carbon”, Sov. Powder Metall. Met. Ceram., 12(36), 1004–1009 (1965), translated from Poroshk. Metall. (Kiev), 12(36), 56–62 (1965) (Experimental, Phase Relations, 5) Rudy, E., Windisch, S., “Part I. Related Binary Systems, Vol. III: Systems Mo-B and W-B”, in “Ternary Phase Equilibria in Transition Metal-B-C-Si Systems” Technical Report AFML-TR-65–2, Part I, Vol. III, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, 72 (1965) (Experimental, Phase Diagram, Phase Relations, 20) Kiparisov, S.S., Nikiforov, O.A., Borisova, N.V., “Some Properties of Alloys in the System W2B-WC” (in Russian), Poroshk. Metall. (Minsk), 353–358 (1966) (Experimental), as cited in [C.A.] Lowell, C.E., “Solid Solution of Boron in Graphite”, J. Am. Ceram. Soc., 50, 142–4 (1967) (Crys. Structure, Experimental, 5) Kiparisov, S.S., Nikiforov, O.A., Yakushin, Yu.S., “Tungsten-Rich Zone of the Tungsten - Boron Carbon System” (in Russian), Sb. Mosk. Inst. Stali Splavov., 45, 121–127 (1968) (Phase Diagram, Experimental), as cited in [C.A.] Kiparisov, S.S., Nikiforov, O.A., Borisova, N.V., “Pseudobinary Section of the Tungsten - Boron Carbon Ternary System” (in Russian), Sb. Mosk. Inst. Stali Splavov, 45, 128–131 (1968) (Experimental), as cited in [C.A.] Rudy, E., “Part V: Compendium of Phase Diagram Data” in “Ternary Phase Equilibria in Transition Metal - Boron - Carbon - Silicon Systems”, Technical Report AFML-TR-65–2, Part V, Air Force Materials Laboratory, Wright Patterson Air Force Base, Ohio, 192–197 (1969) (Experimental, Phase Diagram, Phase Relations, 6) Rudy, E., “Section III.K.4 W-B-C System”, in “Ternary Phase Equilibria in Transition Metal-BoronCarbon-Silicon Systems”, 655–670 (1969) (Crys. Structure, Experimental, Phase Diagram, 3) Rudy, E., “Part V, The Phase Diagram W-B-C” in “Experimental Phase Equilibria of Selected Binary, Ternary and Higher Order Systems”, Report AFML-TR-69–117, Part V, Air Force Materials Laboratory Wright Patterson Air Force Base, Ohio, 1–51 (1970) (Crys. Structure, Phase Diagram, Experimental, #, *, 22) Kuhlmann, H.S., “Determination of the Solubility of Carbon in Tungsten in the Temperature Range from 1400 to 2000˚C” (in German) Tech. Wiss. Abhandl. Osram. Ges., 11, 328–332 (1973) (Phase Diagram, Phase Relations, Experimental), as cited in [C.A.] Harsta, A., Rundquist, S., Thomas, J.O., “A Neutron Powder Diffraction Study of W2C”, Acta Chem. Scand., 32A, 891–892 (1978) (Crys. Structure, Experimental, 10) Schouler, M.C., Ducarroir, M., Bernard, C., “Review on the Constitution and the Properties of the Metal-Carbon-Nitrogen and Metal-Carbon-Boron System” (in French), Rev. Int. Hautes Temp. Refract., 20, 261–311 (1983) (Mechan. Prop., Phase Diagram, Phase Relations, Review, 154) Holleck, H., “Binary and Ternary Carbide and Nitride Systems of the Transition Metals” (in German), Materialkundlich Technische Reihe, Vol. 6, Petzow, G. (Ed.), Gebru¨der Borntra¨ger, Berlin, Stuttgart, 6, 264–274 (1984) (Crys. Structure, Phase Diagram, Phase Relations, Review, 87) Gustafson, P., “Thermodynamic Evaluation of C-W System”, Mater. Sci. Tech., 2, 253–658 (1986) (Phase Diagram, Thermodyn., Review, Phase Relations, 28) Lo¨nnberg, B., Lundstro¨m, T., Tellgren, R., “A Neutron Powder Diffraction Study of Ta2C and W2C”, J. Less-Common Met., 120, 239–245 (1986) (Crys. Structure, Experimental, 17)

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B–C–W [1988Epi]

[1990Ase]

[1991Nag]

[1991Zak]

[1993Wer]

[1993Jon]

[1994Kas] [1994McH]

[1994Par]

[1995Dus]

[1995Oht] [1995Oka]

[1995Vil]

[1996Kas] [1997Bul]

[1998Kas]

[1998Rog1]

[1998Rog2]

24

Epicier, T., Dubois, J., Esnouf, C., Fantozzi, G., Convert, P., “Neutron Powder Diffraction Studies of Transition Metal Hemicarbides M2C1–x - II. In Situ High Temperature Study on W2C1–x and Mo2C1–x”, Acta Metall., 36, 1903–1921 (1988) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 33) Aselage, T.L., Tallant, D.R., Gieske, J.H., “Preparation and Properties of Icosahedral Borides” in “The Physics and Chemistry of Carbides, Nitrides and Borides”, Freer, R. (Ed.), Proc. of the NATO Advanced Research Workshop, Manchester, U.K., Sept. 1989, published as ASI-Series, Series E: Applied Sciences, Vol. 185, Kluwer Acad. Publ., Dordrecht, 97–111 (1990) (Crys. Structure, Review, Experimental, 14) Nagender Naidu, S.V., Rama Rao, P., “C-W (Carbon - Tungsten)” in “Phase Diagrams of Binary Tungsten Alloys”, Nagender Naidu, S.V., Rama Rao, P. (Eds.), Indian Inst. Metals, Calcutta, 37–50 (1991) (Review, Phase Diagram, Phase Relations, 100) Zakhariev, Z., Radev, D., “Dense Material Obtained on the Basis of Boron Carbide Sintered without Pressing” in “AIP Conference Proceedings 231 on Boron-Rich Solids”, Emin, D., Aselage, T.L., Switedick, A.C., Morrosin, B., Beckel, B.C. (Eds.), Albuquerque, USA, (1990), published by AIP, New York, 464–467 (1991) (Experimental, Phys. Prop., 8) Werheit, H., Kuhlmann, U., Laux, M., Lundstro¨m, T., “Structural and Electronic Properties of CarbonDoped β-Rhombohedral Boron”, Phys. Status Solidi B, 179(2), 489–511 (1993) (Crys. Structure, Experimental, 51) Jonsson, S., “An Assessment of the Ti-W-C System”, Trita-Mac 519, The Royal Inst. Technology, Div. Physical Metallurgy, S-10044 Stockholm, (1993) (Phase Diagram, Phase Relations, Thermodyn.), as cited in [C.A.] Kasper, B., Max - Planck - Institute - PML, unpublished work, Stuttgart (1994) McHale, A.E., “VI. Boron Plus Carbon Plus Metal” in “Phase Equilibria Diagrams, Phase Diagrams for Ceramists”, McHale, A.E. (Ed.), Ceramic Div., Natl. Inst. Stand. Technol., Gaithersburg, Maryland, 10, 198–203 (1994) (Phase Diagram, Phase Relations, Review, 4) Parthe, E., Gelato, L., Chabot, B., Penzo, M., Cenzual, K., Gladyshevskii, R., “Typix, Standardized Data and Crystal Chemical Characterization of Inorganic Structure Types”, Vols. 1-4, Gmelin, Handbook of Inorganic and Organometallic Chemistry, Springer, Berlin (1994) (Crys. Structure) Duschanek, H., Rogl, P., “A Critical Assessment and a Thermodynamic Calculation of the Binary System Boron - Tungsten (B-W)”, J. Phase Equilib., 16(2), 150–161 (1995) (Phase Diagram, Phase Relations, Thermodyn., Assessment, #, 50) Ohtani, S., Ohashi, H., Ishizawa, Y., “Lattice Constants and Nonstoichiometry of WB2–x”, J. Alloys Compd., 221, L8-L10 (1995) (Crys. Structure, Experimental, 6) Okada, S., Kudom, K., Lundstro¨m, T., “Preparation and Some Properties of W2B, δ-WB and WB2 Crystals from High-Temperature Metal Solutions”, Jpn. J. Appl. Phys. 34, 226–231 (1995) (Experimental, Crys. Structure, 23) Villars, P., Prince, A., Okamoto, H., “Handbook of Ternary Alloys Phase Diagrams”, Vol. 5, ASM International, Materials Park, Ohio, USA 5357, (1995) (Review, Phase Diagram, Phase Relations, Crys. Structure, 4) Kasper, B., “Phase Equilibria in the B-C-N-Si System” (in German), Thesis, Max-Planck-Institute, Stuttgart, 1–225 (1996) (Calculation, Phase Diagram, Phase Relations, Review, Thermodyn., 170) Bulfon, C., Leithe-Jasper, A., Sassik, H., Rogl, P., “Thermal Expansion and Hardness of Czochralski-grown Single Crystal of WB2–x” in “Symposium 7, Materialwissenschaftliche Grundlagen”, Aldinger, F., Mughrabi, H. (Eds.), DGM-Informationsgesellschaft, 191–196 (1997) (Experimental, Crys. Structure, 8) Kasper, B., Lukas, H.L., “System B-C” in “COST 507. Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), Office for Official Publications of the European Communities, Belgium, Vol. 2, 117–119 (1998) (Review, Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 0) Rogl, P., “The System Boron - Carbon - Tungsten” in “Phase Diagrams of Ternary Metal-Boron-Carbon Systems”, Effenberg, G. (Ed.), MSI, ASM Intl., Materials Park, Ohio, USA, (1998) 372–427 (Experimental, Crys. Structure, Review, Phase Diagram, 32) Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides”, Int. J. Refract. Met. Hard Mater., 17, 27–32 (1999) (Crys. Structure, Thermodyn. Calculation, Phase Relations, Phase Diagram, 6)


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[2000Wen]

[2002Sug] [2004Sug1] [2004Sug2]

[2005Bos] [2005Tan]

[2005Yi]

[2006Pad]

[2006Kar]

[2006Lv1]

[2006Lv2] [2006Lv3] [2006Wen] [2007Lv] [C.A.] [Mas2] [V-C2]

B–C–W Rogl, P., Bittermann, H., “Ternary Metal Boron Carbides - Constitution, Thermodynamics, Compound Formation and Structural Chemistry” in “Materials Science of Carbides, Nitrides and Borides”, Gogotsi, Y., Andreev, R.A. (Eds.), Kluwer Academic Publishers, 29–46 (1999) (Thermodyn., Calculation, Phase Relations, Phase Diagram, 16) Wen, G., Li, S.B., Zhang, B.S., Guo, Z.X., “Processing of in situ Toughened B-W-C Composites by Reaction Hot Pressing of B4C and WC”, Scr. Mater., 43(9), 853–870 (2000) (Experimental, Mechan. Prop., 21) Sugiyama, S., Taimatsu, H., “Preparation of WC-WB-W2B Composites from B4C-W-WC Powders and their Mechanical Properties”, Mater. Trans., 43(5), 1197–1201 (2002) (Experimental, Mechan. Prop., 21) Sugiyama, S., Taimatsu, H., “Preparation of W-C-B Composites by Reactive Resistance-Heated Hot Pressing”, Mater. Sci. Forum, 449-452, pt.1, 309–312 (2004) (Experimental, Mechan. Prop., 16) Sugiyama, S., Taimatsu, H., “Mechanical Properties of WC-WB-W2B Composites Prepared by Reaction Sintering of B4C-W-WC Powders”, J. Eur. Ceram. Soc., 24(5), 871–876 (2004) (Experimental, Mechan. Prop., 24) Bose, K., Wood, R.J.K., “High Velocity Solid Particle Erosion Behaviour of CVD Boron Carbide on Tungsten Carbide”, Wear, 258(1-4), 366–376 (2005) (Experimental, Mechan. Prop., 31) Tanaka, T., Takenouchi, S., Rogl, P., “A Long-Standing Puzzle Solved: Peritectic Reaction L+B4+xC Ð βB”, Research Presented at the 15th International Symposium on Boron, Borides and Related Compounds, Hamburg (Germany), Aug. 21–26, 142 (2005) (Experimental, Phase Relations, 3) Yi, L., Wen, G.W., Song, L.N., Lei, T.Q., Zhou, Y., “Microstructure and Properties of the C/W2B5 Composites Fabricated by Reaction Hot-Pressing”, Rare Metal. Mater. Eng., 34, 381–384 (2005) (Experimental, Morphology, Mechan. Prop., 5) Paderno, Y., Paderno, V., Liashchenko, A., Filipov, V., Evdokimova, A., Martynenko, A., “The Directional Crystallization of W-B-C d-transition Metal Alloys”, J. Solid State Chem., 179, 2939–2943 (2006) (Experimental, Crystal Structure, Phys. Prop., 4) Karaman, M., Sezgi, N.A., Dogu, T., Ozbelge, H.O., “Kinetic Investigation of Chemical Vapor Deposition of B4C on Tungsten Substrate”, Aiche J., 52(12), 4161–4166 (2006) (Experimental, Interface Phenomena, 0) Lv, Y., Wen, G., Zhang, B., Lei, T.Q., “Mechanical Properties and Electrical Conductivity of W-B-C Composites Fabricated by in-situ Reaction”, Mater. Chem. Phys., 97(2-3), 277–282 (2006) (Experimental, Mechan. Prop., Electr. Prop., 19) Lv, Y., Wen, G., Lei, T.Q., “Friction and Wear Behavior of C-based Composites in-situ Reinforced with W2B5”, J. Eur. Ceram. Soc., 26(15), 3477–3486 (2006) (Experimental, Mechan. Prop., 24) Lv, Y., Wen, G., Lei, TQ, “Tribological behavior of W2B5 Particulate Reinforced Carbon Matrix Composites”, Mater. Lett., 60(4), 541–545 (2006) (Experimental, Mechan. Prop., 13) Wen, G., Lv, Y., Lei, T.Q., “Reaction-formed W2B5/C Composites with High Performance”, Carbon, 44(5), 1005–1012 (2006) (Experimental, Mechan. Prop., 23) Lv, Y., Wen, G., Song, L., Lei, T.Q., “Wear Performance of C-W2B5 Composite Sliding Against Bearing Steel”, Wear, 262(5-6), 592–599 (2007) (Experimental, Mechan. Prop., 17) Chemical Abstracts - pathways to published research in the world’s journal and patent literature - http:// www.cas.org/ Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)

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Boron – Carbon – Zirconium Refractory Metal Systems: Phase Diagrams, Crystallographic and Thermodynamic Data Peter Rogl

Introduction Ultra-high-temperature ceramics, composites and cermets, ZrB2/ZrC1–x and ZrB2/B4C have attracted significant attention because of a unique combination of materials properties such as high thermal conductivity, good thermal shock resistance at moderate thermal expansion and oxidation resistance. Based on early studies of the interaction between zirconium borides, zirconium carbide and carbon [1952Gla, 1955Bre, 1959Sam], phase relations in the B-C-Zr ternary system have been established by [1960Now, 1961Now], who investigated the isothermal section at 1400˚C, and by [1965Lev], who provided preliminary information on the quasibinary eutectic ZrB2-(C)gr. The most complete experimental information, however, is from [1966Rud], comprising a reinvestigation of the isothermal section at 1400˚C, a partial isothermal section for the B rich region at 1900˚C, the determination of the liquidus surface and the investigation of three isopleths, i.e. ZrB2-ZrC0.88, ZrB2-C and ZrB2-B4.5C. On the basis of these experimental results, [1966Rud] derived a reaction scheme for the ternary system and constructed a series of further isotherms at 1800, 2160, 2300, 2400, 2600, 2800 and 3000˚C, as well as a threedimensional isometric view of the ternary system. More recent and independent studies of the isopleths ZrB2-ZrC1–x and ZrB2-‘B4C’ confirmed the quasibinary and eutectic nature of these sections [1975Ord, 1988Ord, 2000Kov]. Whilst the eutectic temperature TE = 2280±30˚C of [1988Ord] for ZrB2-‘B4C’ is in acceptable agreement with the data given by [1966Rud] (TE = 2220±20˚C), the concentration of the eutectic point at 75 mol% ‘B4C’ [1988Ord], or 71.5 mol % ‘B4C’ [2000Kov] is considerably richer in ‘B4C’ than recorded by [1966Rud] (65±5 mol% B4.5C). Besides the redetermination of the ZrB2-ZrC1–x quasibinary eutectic section (TE = 2660±40˚C at 43 mol% ZrC1–x) by [1975Ord], the studies of [1979Tka, 1979Fri, 1982Unr] are mainly concerned with the physical properties in the system ZrB2-ZrC1–x and particularly with the rather low microhardness for the eutectic composition. In this case the eutectic composition is close to the value of 42 mol% Zr0.54C0.46 given by [1966Rud] but the eutectic temperature seems to be rather low when compared to [1966Rud] (2830±15˚C). Compilations of the most prominent features of the B-C-Zr phase diagram were presented by [1969Rud, 1974Upa, 1983Sch, 1984Hol, 1991Ruy, 1994McH, 1995Vil]. A full status of all information in literature for the B-C-Zr system up to 1996 was compiled in a general review of phase relations for metal-boron-carbon systems [1998Dus] including a thermodynamic calculation of the entire B-C-Zr system. For further details on the thermodynamic assessment, see [1999Rog]. Experimental details for all investigations in the B-C-Zr system are summarized in Table 1.


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Binary Systems The C-Zr system is taken from a critical assessment and thermodynamic calculation by [1995Gui] (Fig. 1). A partial low temperature phase diagram C-Zr was calculated by [1991Gus] integrating the formation of ordered superlattice phases ‘Zr2C’, ‘Zr3C2’, ‘Zr6C5’, all deriving from the NaCl type parent phase ZrC1–x. These phases were calculated to form below about 900˚C. A critical assessment and thermodynamic calculation of the B-Zr binary system is from [1988Rog] and in a refined version from [1998Dus] (Fig. 2) including a discussion on the (αZr)/(βZr) transition as well as on the temperature region of existence of ZrB12. The B-C system is adopted from a recent assessment and thermodynamic calculation by [1996Kas, 1998Kas]. The phase diagram is included in the present volume in the evaluation of the B-C-Cr system. The boron rich reaction, L+‘B4C’ Ð (βB), as modelled by [1996Kas], was experimentally confirmed from floating zone experiments on several carbon-doped boron samples [2005Tan]. Crystallographic data for the binary boundary phases are summarized in Table 2.

Solid Phases No ternary B-C-Zr compounds have been reported. Mutual solid solubilities among the binary borides and carbide phases generally were found to be very small [1960Now, 1961Now, 1965Lev, 1966Rud, 1975Ord, 1982Unr, 1988Ord, 2000Kov], except for the zirconium monocarbide, for which lattice parameters in the ternary are considerably increased with respect to those of the binary [1966Rud] (Table 2). The ternary solid solubility of the nonmetal elements in Zr at the temperature of the ternary eutectic E2 (L Ð (βZr) + ZrB2 + ZrC1-x) was said to be smaller than 1 at.% altogether and a peritectoid decomposition of the (αZr) phase on heating was suggested [1966Rud] (see also section “Invariant Equilibria”). ZrB12 was absent in the ternary alloy series annealed at 1400˚C and 1600˚C; moreover ternary alloys annealed at 1900˚C revealed ZrB12 with a lattice parameter practically identical with the binary value of a = 740.8 pm from which a small or negligible homogeneity region was concluded [1966Rud].

Quasibinary Systems Three quasibinary systems of the eutectic type were established by [1966Rud]. The eutectic nature of the ZrB2-C quasibinary was earlier claimed by [1965Lev]. The eutectic temperature, 2227±30˚C as measured by optical pyrometry on pre-reacted powders through a bore hole in a directly heated graphite tube, is, however, remarkably low compared to 2390±15˚C reported by [1966Rud] and probably explains from insufficient correction for non-blackbody conditions in the experiment of [1965Lev]. Furthermore, the eutectic composition at 81 mol% C [1965Lev] is in distinct disagreement with the eutectic at 33±2 mol% C [1966Rud]. Recrystallization of the ZrB2-ZrC1-x quasibinary eutectic at temperatures close to the eutectic line was said to be extremely fast and partially or fully annealed structures resulted, if cooling rates lower than about 30 K·s–1 were employed [1966Rud]. Alloys crossing the homogeneous range of the monocarbide at 5 at.% B were found to be single phase when rapidly cooled from 2900 to 3400˚C; diboride precipitation from the zirconium monocarbide, however, was observed to DOI: 10.1007/978-3-540-88053-0_25 ß Springer 2009

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be much slower than in the corresponding Ti-containing alloys [1966Rud]. The absence of interaction between ZrC1–x and ZrB2 at T ≤ 2100˚C was confirmed from powder X-ray and microhardness investigations on samples compacted from Zr + C + ZrB2 powders presintered at 2100˚C and heated at temperatures above 2300˚C [1975Ord, 1979Fri, 1979Tka, 1982Unr]. Solid solubility of ZrB2 in ZrC1–x in samples quenched from 2400˚C was said to be about 3 mol% ZrB2 and was claimed [1975Ord] to remain virtually identical after reheating to 2100˚C, as seen from the practically unchanged lattice parameter of Zr(C,B)1–x (a = 470.0 pm). Whereas the eutectic composition at 43 mol% ZrC1–x, reported by [1975Ord] is in close agreement with that of [1966Rud] (42 ±⊊2 mol% Zr0.54C0.46), the eutectic temperature recorded at 2660±40˚C [1975Ord] is relatively low, when compared to 2830±15˚C [1966Rud] and may be explained by insufficient correction for nonblack body conditions in pyrometric recording. From the carbon content of the ZrC1–x powder used by [1975Ord], the actual concentration of the section investigated is ZrB2 - ZrC0.95. The lattice parameter of ZrC1–x (a = 470.0 pm), however, will point towards the carbon poor side of the Zr(C,B)1–x solution, where solubility of ZrB2 according to [1966Rud] was observed to be much larger, i.e. 5 mol% Zr0.33B0.67. It shall be noted, that taking the aforementioned sections through the congruent melting points of the binary compounds does not constitute true quasibinary sections as seen for instance from the slight deviation in the sections ZrB2-(C)gr (see Fig. 3) or ZrB2-ZrC1–x (see Fig. 4). Deviations are, however, small and experimentally inaccessible. Acceptable agreement exists on the eutectic nature, the eutectic temperature and to a lesser extent on the eutectic composition in the section ZrB2-B4.5C (Fig. 5) investigated by [1966Rud] (TE = 2220±20˚C, 65 mol% B0.817C0.183) and by [1988Ord] (TE = 2280"30˚C, 75 mol% B4.5C). From a directionally solidified eutectic alloy (zone melted) [2000Kov] determined the eutectic composition from chemical analysis to be 71.5 mol% ‘B4C’ (in at.% Zr6.4B77.4C16.2). The eutectic structure obtained was fibrous ZrB2 embedded in a continuous matrix of ‘B4C’ [1966Rud, 1988Ord, 2000Kov].

Invariant Equilibria A Scheil reaction scheme including nine observed ternary invariant equilibria was provided by [1966Rud]. Table 3 lists the compositions of the phases at the four-phase isothermal reactions plus the maximum points in the liquidus trough as given by [1966Rud] and compares experimentally derived data with the results of the thermodynamic calculation by [1998Dus, 1999Rog]. Figure 6 shows the Scheil diagram corresponding to the thermodynamic assessment. [1966Rud] gave the (βZr)/(αZr) transformation as a mean value of eight individual observations in the B-C-Zr ternary system at 880±15˚C suggesting a ternary peritectoid formation of (αZr) (calculated equilibrium P4 at 923˚C; see Fig. 6). As for the ternary eutectic L Ð (βZr) + ZrC1–x + ZrB2 at 1615˚C (calculated equilibrium E2 at 1652˚C; see Fig. 6), annealing reactions close to the eutectic temperature were said to be very fast and with respect to the clustering tendency of ZrB2, quenching rates >150 K·s–1 were essential to retain the structures in the as crystallized state. As far as the very boron rich region is concerned, the thermodynamic calculation renders a transition type reaction U3 at 2018˚C, L + ‘B4C’ Ð ZrB12+ (βB), rather than a ternary eutectic at 1990˚C, L Ð (βB) + ‘B4C’ + ZrB12, as reported by [1966Rud]. This discrepancy essentially results from a new assessment of the B-C system by [1996Kas] revealing a peritectic reaction Landolt‐Bo¨rnstein New Series IV/11E1

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L + ‘B4C’ Ð (βB) at 2103˚C rather than a eutectic L Ð (βB) + ‘B4C’ at 2080˚C as given by [1966Rud]. The experimentally reported very sharp drop of almost 100˚C and within less than 2 at.% from the binary reaction isotherms into the ternary eutectic (at 1990˚C [1966Rud]) seems unlikely and could not be modeled thermodynamically. It shall be noted that the peritectic binary reaction L + ‘B4C’ Ð (βB) was recently confirmed from floating zone measurements on several carbon-doped boron samples [2005Tan]. The slight deviation i.e. incompatibility with a true quasibinary for the calculated section ZrB2+C shows further consequences insofar as the experimentally observed ternary eutectic, L Ð ZrB2 + ZrC1–x + (C)gr at 2360˚C [1966Rud] turns into a transition reaction U1, L + ZrC1–x Ð ZrB2 + (C)gr calculated at 2369˚C [1998Dus, 1999Rog].

Liquidus, Solidus and Solvus Surfaces Figure 7 represents the liquidus surface projection calculated at intervals of 200 K. Figure 8 is a projection of the liquidus troughs in the Gibbs triangle as a function of temperature and as seen along the B-Zr axis in direction of boron.

Isothermal Sections Based on the experimentally determined isothermal sections at 1400, 1900˚C as well as on the liquidus projection and from the phase relations experimentally established for the three isopleths ZrB2-ZrC1–x, ZrB2-C and ZrB2-B4.5C, [1966Rud] constructed a series of isothermal sections at 1800, 2160, 2300, 2400, 2600, 2800 and 3000˚C. A set of sections, calculated by [1998Dus], is shown in Figs. 9 to 14.

Temperature – Composition Sections Isopleths ZrC-B and Zr-B0.5C0.5 were derived by [1966Rud] from the experimental data comprising isothermal sections at 1400, 1900˚C, the liquidus projection and the quasibinary sections ZrB2-C, ZrB2-ZrC1–x and ZrB2-B4.5C. In order to elucidate phase interactions in the ternary system a series of isopleths was calculated [1998Dus, 1999Rog]: ZrC1–x- B (Fig.15), Zr - B0.817C0.183 (Fig. 16) and ZrB12 - B0.817C0.183 (Fig. 17). Small deviations from congruent melting of the constituents can significantly change the extent of the phase fields as seen for the sections ZrC-‘B4C’ (Fig. 18) and ZrB2-ZrC (Fig. 19) calculated for ZrC at the C rich phase boundary.

Thermodynamics A thermodynamic calculation of the ternary B-C-Zr system by [1998Dus, 1999Rog] was based on thermodynamic assessments of the binary systems B-C [1996Kas], B-Zr [1998Dus] and C-Zr [1995Gui] as well as relying on the phase diagram data from [1966Rud] for the optimization of the thermodynamic parameters. DOI: 10.1007/978-3-540-88053-0_25 ß Springer 2009

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Notes on Materials Properties and Applications The combination of refractory metal like borides with high temperature ceramics constitutes attractive high temperature materials. Therefore considerable interest is devoted to the manufacture, high temperature behavior and properties of ZrB2/ZrC and ZrB2/‘B4C’ composites or ZrB2 based materials in general. Starting from the improvement of sinterability of ‘B4C’ on additions of Zr or ZrN [1978Bod], many articles deal with various techniques on pressureless sintering of dense ZrB2/B4C composites [2001Gol, 2004Tsu1, 2004Tsu2, 2004Tsu3, 2006Liu1], ZrB2-toughened ‘B4C’ [2006Liu2], structural changes in ‘B4C’ by Zr introduction [1999Liu, 2000Liu], pressureless densification of ZrB2 by ‘B4C’- additions [2006Zha], pressureless sintering of carboncoated zirconium diboride powders [2007Zhu], spark plasma sintering of ZrB2/ZrC [2007Tsu] and oxidation behavior of ZrB2-‘B4C’ composites [1993Rad]. Microstructure and kinetics of the formation of a new class of ZrB2 platelet reinforced ZrC ceramic composites were studied on material made by strongly exothermic reaction of molten Zr with B4C yielding ZrB2-platelets in a ZrC matrix with a controlled amount of residual Zr [1989Cla, 1989Whi, 1992Bre, 1991Joh]. [2007Li] reported on nanoscale ZrB2/ZrC multilayered coatings prepared by magnetron sputtering reaching a hardness of 47 GPa. Strength and deformation properties, thermal conductivity and diffusivity, friction and wear, were investigated by [1959Sam, 1979Fri, 1979Tka] on quasibinary ZrB2-ZrC1–x eutectic alloys with a grain size of 3 mm. Eutectic alloys were found to reveal minimal wear, maximal friction, maximal bending and compressive strengths and minimal hardness (13.2 to 19.6 GPa) when compared to noneutectic ZrB2-ZrC1–x compositions [1975Ord, 1979Tka] and are characterized by the absence of marked grain growth. Cracks propagated along the grain boundary. In hypereutectic alloys ZrB2 grows preferentially along the c-axis in form of long needles and lamellae [1975Ord]. Dependency of the morphology of the ZrB2-ZrC1–x eutectic as a function of cooling rates was studied by [1982Unr]. The growth direction of the eutectic is {01–1}ZrC and {010}ZrB2; accordingly the morphology of directionally solidified eutectic ZrB2-ZrC consists of columnar grains of parallel lamellae with epitaxial relationship {111}ZrC// {001}ZrB2 (growth law λ2R = (2.57±0.63) 10–18 m3·s–1; lattice parameter mismatch 4.6%) [1984Sor] (note, that a most regular eutectic growth was obtained for 48 mol% ZrB2 - the compositional discrepancy was explained by the occurrence of banding). [2002Shi] investigated the crystallographic orientation of ZrB2-ZrC composites manufactured by spark plasma sintering. Microhardness measurements on samples in both the isopleths ZrB2 - ZrC1–x and ZrB2 - ‘B4C’ near the eutectic composition were found to reveal values well below the linear combination of the binary compounds and to be strongly dependent on the eutectic crystallization conditions. For the ZrB2 - ZrC1–x eutectic, microhardness ranges from 13.2 GPa for the finely dispersed eutectic to 19.6 GPa for macrocrystalline colonies [1975Ord]. For the ZrB2 ‘B4C’ eutectic microhardness was reported to be 32 to 33 GPa [1988Ord, 2000Kov]. Dependency of the morphology of the ZrB2-‘B4C’ eutectic as a function of cooling rates was studied by [2000Kov].

Miscellaneous [1968Alp] discussed phase assemblages, microstructures, and properties of fused carbides and/or borides from refractory B-C-M systems (also for B-C-Zr) containing free graphite in Landolt‐Bo¨rnstein New Series IV/11E1

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terms of compatibility and phase diagrams. All these bodies have excellent thermal-shock resistance. Other properties (such as electrical, thermal, mechanical, chemical) can be modified by choosing different phase assemblages. Some of these materials have been cast into large shapes more than 45 cm long, and they can be machined into articles. Interaction in the quasibinary eutectic systems MC - MB2 [1980Ord] and MB2 - ‘B4C’ [1993Ord] has been analyzed for transition elements M = Ti, Zr, Hf, V, Nb and Ta and correlations were found between the relative quasieutectic temperature and the d5 electron concentration of the metal atoms. ZrB2 + ‘B4C’ + SiC ceramics showed low specific wear rate at 800˚C in air [1995Ume]. Properties of B-C-Zr alloys made by pyrolytic decomposition of ZrCl4-BCl3-natural gas mixtures at 1300–2100˚C were described by the authors of [1969Der]. C5H5-Zr(BH4)3 has proven to be a good precursor for the plasma enhanced CVD of Zr(C,B) films [1994Rei]. Zirconium and boron containing MO-PACVD coatings were developed for the application on aluminum die casting tools [1997Rie]. A sintered ZrB2 - 3–30% ‘B4C’ body may be used as crucible for Czochralski or Bridgman growth of tantalates or niobates (KNbO3) [1986Sun]. Using a compact diffraction reaction chamber, [2006Won] studied with time resolved X-ray diffraction the chemical dynamics at the combustion front for the directly ignited reactions Zr+C = ZrC, Zr+2B = ZrB2 and 3Zr+B4C = 2ZrB2+ZrC. Combustion front velocities were 5.8, 5, and 6.4 mm·s–1, respectively. The adiabatic combustion temperatures calculated (>3000 K) exceeded the Tm of Zr in all reactions.

. Table 1 Investigations of the B-C-Zr Phase Relations, Structures and Thermodynamics Temperature/Composition/Phase Range Studied

Reference


[1952Gla]

Hot-pressing of powder mixtures (powders of ZrH2, ZrC and ZrB2,) in graphite dies at 1300 to 2400˚C. XPD on six hot-pressed samples.

[1955Bre]

Reaction between metal borides and XPD on a sample (Zr + 2B + C). graphite. Sintering of powder compacts in Mo-crucibles under 0.5 bar argon for 50 min at 1777˚C.

Compatibility of ZrB2 with C and with ‘B4C’. Small amounts of C in ZrB2+B form ZrB12.

Investigation of the sections ZrC-ZrB, [1959Sam] Samples were prepared from ZrO2 via borothermic reaction with B4C or B2O3+C. ZrB2-ZrC. Continuous solid solution of Zr LOM, XPD and microhardness. (C1–xBx), immiscibility of ZrB2-ZrC. Oxidation in air of ZrB2+ZrC, as well as friction and wear were investigated.

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[1960Now] Hot-pressing of 100 binary and ternary Isothermal section at 1400˚C. Study of the [1961Now] powder compacts in C-cartridges at 1400 homogeneity region of Zr(C,B)1–x at 1400˚ to 2500˚C. A series of alloys was prepared C from lattice parameter dependencies. by reaction sintering in Ar or vacuum respectively. Samples containing B, B4C, C, ZrC and ZrB2 were reacted at 2000˚C for several h prior to final anneal at 1400˚C for 3 to 16 h. Alloys containing free Zr were annealed for 4 h at 1400˚C. Starting materials were ZrH2 containing 0.4% O), amorphous boron (96 mass% residue O, C), lampblack C. Master alloy ZrC (12.33 mass% Ctotal, 1.46 mass% free C). Some alloys were arc melted. XPD and LOM. [1965Lev]

ZrB2 was prepared via the boron carbide process at 2000 to 2200 K, MO2+‘B4C’+C = MB2+CO. Starting materials were powders of ZrO2, ‘B4C’ (containing B2O3, C), lampblack C (ash 4˚/sec) techniques. Selected alloys were equilibrated in the melting point furnace and quenched in a preheated tin bath (300˚C). Samples were chemically analyzed for free and combined carbon, boron, oxygen and nitrogen contaminants. For polishing and etching usually a slurry of alumina in 5% chromic acid was used; for alloys within the nominal compositions Zr-Zr0.6C0.4Zr0.4B0.6 anodic oxidation in an electroetching process using 10% oxalic acid was said to provide excellent phase contrast coloring the metal phase light blue, the monoboride brown, whereas the diboride remains unaffected. Specimens from the region Zr-Zr0.6C0.4Zr0.4B0.6-Zr0.3B0.7-Zr0.2C0.8 were dip-etched in a solution with 10% aqua regia and HF.


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. Table 1 (continued) Temperature/Composition/Phase Range Studied

Reference


[1975Ord]

Starting materials were ZrC (88.2 mass% Zr, Ctot = 11.2 mass% C, 0.15 mass% free C) and ZrB2 (80.1 mass% Zr, Btot = 18.11 mass%, 0.06 mass% free C, 0.7 mass % (O2+N2)). Specimens were obtained from extrusion of plasticized masses in form of cylinders (3 mm diameter · 50 mm length), presintered at 2100˚C for 4 h in vacuum prior to heat treatment at 2300˚C under Ar. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

Investigation of the pseudo-eutectic system ZrB2+ZrC1–x with eutectic point at 57 mol% ZrB2 and TE of 2660±40˚C on 9 samples.

[1984Sor]

Preparation of 35 g sinter bars (1·1·10 cm3) from 325 mesh powders of 98.6 mass% pure ZrB2 (0.6 mass% C, 0.3% N, 0.19% O plus Fe, H) and 0.99.8% ZrC0.90 (containing Fe, Sn, O, Al, Cu) at 2300˚C for 15 min under Ar. LOM, TEM.

Determination of ZrB2+ZrC1–x eutectic morphology on floating zone melted rod by LOM, TEM. A most regular eutectic growth was obtained for 48 mol% ZrB2.

[1988Ord]

Investigation of the quasieutectic system Samples were prepared from powder ZrB2+‘B4C’ with eutectic point at 25 mol% compacts of ZrB2 and ‘B4C’ (which was vacuum annealed at 2000˚C to reduce C ZrB2 and TE of 2280±30˚C on 10 samples. content to 0.2 mass% free C). Specimens in form of cylinders (3 mm diameter · 50 mm length) were compacted with aid of 12% aqueous starch solution, presintered at 2300˚C for 2 h in vacuum prior to measurement. Specimens with high ‘B4C’ content were indirectly melted inside a W-spiral furnace. Chemical analyses, LOM, XPD, microhardness. Pirani-Alterthum melting point analyses under argon.

[2000Kov]

Floating zone refinement under 4.9 104 Pa Determination of the eutectic point in the Ar on rounded off ingots (10·10·100 mm3) quasieutectic system ZrB2+‘B4C’. with starting compositions (in mol%) 70, 75, 80 ‘B4C’. The composition of the eutectic was obtained via chemical analysis. Commercial powders ZrB2, ‘B4C’ were vacuum annealed at 1950˚C for 1 h to reduce Cfree to 65 65

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